ceramic-matrix composites microstructure

i

Ceramic matrix composites

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ii

Ceramic matrixcomposites

Microstructure, properties andapplications

Edited by

I. M. Low

Woodhead Publishing and Maney Publishingon behalf of

The Institute of Materials, Minerals & Mining

CRC PressBoca Raton Boston New York Washington, DC

W O O D H E A D P U B L I S H I N G L I M I T E DCambridge England

iii

Woodhead Publishing Limited and Maney Publishing Limited on behalf ofThe Institute of Materials, Minerals & Mining

Woodhead Publishing Limited, Abington Hall, Abington,Cambridge CB1 6AH, Englandwww.woodheadpublishing.com

Published in North America by CRC Press LLC, 6000 Broken Sound Parkway, NW,Suite 300, Boca Raton, FL 33487, USA

First published 2006, Woodhead Publishing Limited and CRC Press LLC© Woodhead Publishing Limited, 2006The authors have asserted their moral rights.

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iv

Contents

Contributor contact details

Introduction 1

Part I: Fibre-whisker- and particulate-reinforcedceramic composites

1 Fibrous monolithic ceramics 9Y-H KOH, Seoul National University, Korea

1.1 Introduction 91.2 History 91.3 Processing 111.4 Structures 141.5 Mechanical properties 151.6 Future trends 281.7 References 29

2 Whisker-reinforced silicon nitride ceramics 33M D PUGH, Concordia University, Canada and M BROCHU,McGill Univeristy, Canada

2.1 Introduction 332.2 Fabrication 342.3 Properties 372.4 Applications 542.5 References 55

3 Fibre-reinforced glass/glass-ceramic matrixcomposites 58R BANERJEE and N R BOSE, Central Glass and CeramicResearch Institute, India

3.1 Introduction 583.2 Types of fibre suitable as reinforcements in

different glass/glass-ceramic matrix composites 60

v

3.3 Methods for manufacturing different fibre-reinforcedglass/glass-ceramic matrix composites 72

3.4 Properties of glass/glass–ceramic matrix composites 803.5 Microstructural observation 923.6 Application areas 933.7 Future trends 953.8 References 95

4 Particulate composites 99R I TODD, University of Oxford, UK

4.1 Introduction 994.2 Powder processing and microstructural development 1004.3 Thermal microstresses 1034.4 Toughening 1054.5 Room-temperature strength 1104.6 High-temperature strength 1164.7 Wear 1204.8 Future trends 1244.9 References 125

Part II Graded and layered composites

5 Functionally-graded ceramic composites 131I M LOW, R D SKALA and P MANURUNG, Curtin University ofTechnology, Australia

5.1 Introduction 1315.2 Infiltration kinetics and characteristics 1325.3 Infiltration processing of LGMs 1375.4 Characterisation and properties of alumina-matrix LGMs 1385.5 Concluding remarks 1505.6 Acknowledgements 1515.7 References 151

6 SiAION based functionally graded materials 154H MANDAL, Anadolu University, Turkey and N ÇALIS ACIKBAS,MDA Advanced Ceramics, Ltd, Turkey

6.1 Introduction 1546.2 Functionally graded materials 1546.3 SiAlON ceramics 1556.4 Functionally graded SiAlON ceramics 1606.5 Production techniques of functionally graded SiAlON

ceramics 1616.6 Concluding remarks 1746.7 References 175

Contentsvi

7 Design of tough ceramic laminates by residualstresses control 178N ORLOVSKAYA, Drexel University, USA, M LUGOVY, Institute forProblems of Materials Science, Ukraine, J KUEBLER, EMPA, Labfor High Performance Ceramics, Switzerland, S YARMOLENKO andJ SANKAR, North Carolina A&T State University, USA

7.1 Introduction 1787.2 Laminate design for enhanced fracture toughness 1797.3 Processing of Si3-T4–TiN and B4C–SiC ceramic laminates 1897.4 Si3N4 based laminates 1937.5 B4C based laminates 2017.6 Future trends 2107.7 Acknowledgements 2117.8 References 211

8 Hardness of multilayered ceramics 216W J CLEGG, F GIULIANI, Y LONG, S J LLOYD, University ofCambridge, UK and J M MOLINA-ALDAREGUIA, Centro deEstudios e Investigaciones Tecnicas de Gipuzkoa (CEIT), Spain

8.1 Introduction 2168.2 Behaviour of multilayer structures 2178.3 Hardening mechanisms in multilayers 2198.4 Microstructural changes due to making a multilayer 2308.5 Conclusions 2368.6 Future trends 2378.7 Further reading 2378.8 References 237

Part III: Nanostructured ceramic composites

9 Nanophase ceramic composites 243L YONGLI, Beijing University of Technology, China

9.1 Introduction 2439.2 Micro–nano type ceramic composites 2449.3 Nano–nano type ceramic composites 2489.4 Fabrication of nanoceramics 2559.5 Conclusions and future trends 2579.6 References 257

10 Nanostructured coatings on advanced carbonmaterials 260Y MORISADA and Y MIYAMOTO, Osaka University, Japan

10.1 Introduction 26010.2 Coating method of nanostructured SiC 261

Contents vii

10.3 Applications of nanostructured SiC coatings in advancedcomposites 273

10.4 Conclusions 28110.5 References 281

11 Processing and microstructural control of metal-reinforced ceramic matrix nanocomposites 285W D KAPLAN, and A AVISHAI, Technion – Israel Institute ofTechnology, Israel

11.1 Introduction 28511.2 Processing 28511.3 Microstructure 29011.4 Properties 30011.5 Future trends 30411.6 References 304

12 Carbon nanotubes-ceramic composites 309A PEIGNEY and CH LAURENT, CIRIMAT, Université Paul-Sabatier,France

12.1 Introduction 30912.2 Structure, synthesis and properties of carbon nanotubes 30912.3 Preparation of CNT-ceramic composites 31312.4 Properties of CNT-ceramic composites 32012.5 Conclusions and future trends 32912.6 Sources of further information 33012.7 References 331

13 Machinable nanocomposite ceramics 334R WANG, Arizona State University, USA

13.1 Introduction 33413.2 Design principles of machinable ceramics 33413.3 Al2O3–LaPO4 33513.4 Si3N4/h-BN 34413.5 Machinable nanocomposites 35113.6 Conclusions 35213.7 References 354

Part IV: Refractory and speciality ceramic composites

14 Magnesia–spinel (MgAl2O4) refractory ceramiccomposites 359C AKSEL, Anadolu University, Turkey and F L RILEY, Universityof Leeds, UK

14.1 Introduction 359

Contentsviii

14.2 Crystal structures 36214.3 Production of MgAl2O4 36214.4 Densification 36314.5 In-situ formed/preformed spinel based refractories 36514.6 Industrial applications and properties of magnesia–spinel

materials 36614.7 Thermal shock 37214.8 Mechanical properties and thermal shock behaviour of

magnesia–spinel composite refractory materials 37514.9 Conclusions 38914.10 Future trends 39014.11 Acknowledgements 39214.12 Sources of further information 39214.13 References 393

15 Thermal shock of ceramic matrix composites 400C KASTRITSEAS, P SMITH and J YEOMANS, University of Surrey, UK

15.1 Introduction 40015.2 Thermal shock of brittle materials: the induced stress field 40115.3 Experimental methods 40715.4 Thermal shock of monolithic ceramics 41015.5 Thermal shock of particle- and whisker-reinforced CMCs 41315.6 Thermal shock of fibre-reinforced CMCs 41615.7 Concluding remarks 42715.8 References 428

16 Superplastic ceramic composites 434A DOMÍNGUEZ-RODRÍGUEZ, D GÓMEZ-GARCÍA, Universidad de Sevilla,Spain, and F WAKAI, Tokyo Institute of Technology, Japan

16.1 Introduction 43416.2 Macro- and microscopic superplastic characteristics 43516.3 Accommodation processes controlling superplasticity 43916.4 Parameters improving superplasticity 44516.5 Applications of superplasticity 44816.6 Future trends 45216.7 Acknowledgements 45416.8 References 454

Part V: Non-oxide ceramic composites

17 Interfaces in non-oxide ceramic composites 461S TURAN, Anadolu University, Turkey and K M KNOWLES,University of Cambridge, UK

17.1 Introduction 461

Contents ix

17.2 Assessment of the accuracy of TEM techniques for thedetection and measurement of film thickness at interfaces 463

17.3 Wetting, non-wetting and dewetting behaviour ofinterphase boundaries in non-oxide ceramic composites 467

17.4 Equilibrium film thickness at interphase boundaries 46917.5 Effect of intergranular film composition on equilibrium

film thickness 47317.6 Crystallography of interphase boundaries 47517.7 Future trends 48217.8 Further reading 48417.9 References 486

18 Sialons 491Z B YU, Queen’s University, Canada and D P THOMPSON,University of Newcastle, UK

18.1 Introduction 49118.2 Sialons 49218.3 Challenges in toughening and strengthening sialons 49318.4 Sialon composites 49418.5 Future trends 51018.6 References 510

19 Carbon-ceramic alloys 514C BALÁZSI, Research Institute of Technical Physics andMaterials Science, Hungary

19.1 Introduction 51419.2 Carbon as fugitive additive for porous silicon nitride

processing 51519.3 Comparison of silicon nitrides with carbon additions

prepared by hot isostatic pressing and pressureless sintering 51819.4 In situ processing of Si3N4/SiC composites by carbon

addition 52419.5 Silicon nitride ceramics reinforced with carbon fibres and

carbon nanotubes 53019.6 Concluding remarks 53319.7 References 533

20 Silicon nitride and silicon carbide-based ceramics 536Y ZHANG, New York University College of Dentistry, USA

20.1 Introduction 53620.2 Material selection 53720.3 Material characterization 54020.4 Erosion response 54320.5 Microstructure and mechanical properties 550

Contentsx

20.6 Microstructure and erosion mechanisms 55220.7 Conclusions 55620.8 References 557

21 Oxynitride glasses – glass ceramics composites 560S HAMPSHIRE, University of Limerick, Ireland

21.1 Introduction 56021.2 Potential applications 56021.3 Oxynitride glass/glass ceramic composites 56121.4 Oxynitride glass–silicon carbide composites 56921.5 Conclusion 57221.6 References 572

22 Functionally graded ceramics 575G ANNÉ, J VLEUGELS and O vAN DER BIEST, KatholiekeUniversity Leuven, Belgium

22.1 Introduction 57522.2 Functionally graded ceramics concept 57522.3 Classifications of FG ceramics 57722.4 Processing of FGMs 57722.5 FGM design for structural applications 58422.6 Future trends 59022.7 Further reading 59122.8 References 591

Index 597

Contents xi

Contributors contact details

(* = main contact)

Chapter 1

Professor Young-Hag KohSchool of Materials Science andEngineeringSeoul National UniversitySeoul, 151-742Korea

Tel: 82-2-880-1397Fax: 82-2-884-1413E-mail: [email protected]

Chapter 2

Dr Martin Pugh*Department of Mechanical andIndustrial EngineeringConcordia University1455 de Maisonneuve Blvd, WestMontréal, Québec, H3G 1M8Canada

E-mail: [email protected]

Dr Mathieu BrochuDepartment of Mining, Metals andMaterials Engineering3610 University StreetMontréal, Québec H3A 2B2Canada


Chapter 3

Dr Rajat Banerjee* andDr Nripati Ranjan BoseCentral Glass and CeramicResearch InstituteJadavpurKolkata-700 032India


Chapter 4

Dr Richard I. ToddUniversity of OxfordDepartment of MaterialsParks RoadOxfordOX1 3PHUK

E-mail:[email protected]

Chapter 5

Professor It-Meng (Jim) Low*Department of Applied PhysicsCurtin University of TechnologyGPO Box U1987Perth, WA 6845Australia


xiii

Contributor contact detailsxiv

Dr Robert D. SkalaMillennium Chemicals, Inc.6752 Baymeadow DriveGlen BurnieMaryland, MD 21060USA

Dr P. ManurungDepartment of PhysicsUniversity of LampungBandar Lampung35145Indonesia

Chapter 6

Professor Dr Hasan Mandal*Anadolu UniversityDepartment of Materials Scienceand Engineering26470, EskisehirTurkey

Tel: +90 222 322 36 62Fax: +90 222 323 95 01E-mail: [email protected]

Nurcan Çalıs AcıkbasMDA Advanced Ceramics LtdOrganize Sanayii BölgesiEskisehirTurkeyTel: +90 222 2301880Fax: +90 222 2301881E-mail: [email protected]

Chapter 7

Professor N. Orlovskaya*Department of Materials Scienceand EngineeringDrexel University3141 Chestnut StreetPhiladelphia, PA 19104USA


Dr M. LugovyInstitute for Problems of MaterialsScience3 Krzhizhanovskii Street03142 KievUkraine


J. KueblerEMPA, Laboratory for HighPerformance CeramicsUeberlandstrasse 129CH-8600 DuebendorfSwitzerland


S. Yarmolenko and J. SankarDepartment of MechanicalEngineeringNorth Carolina A&T StateUniversity1407 E. Market St.Greensboro, NC 27411USA

E-mail: [email protected]: [email protected]

Chapter 8

Dr William John Clegg*,F Giuliani, Dr Y. Long,Dr S. J. LloydGordon LaboratoryDepartment of Materials Scienceand MetallurgyUniversity of CambridgePembroke StreetCambridge CB2 3QZUK

Tel: +44 (0)1223 334470E-mail: [email protected]

Dr J. M. Molina-AldareguiaCentro de Estudios eInvestigaciones Tecnicas deGipuzkoa (CEIT)Paseo de Manuel Lardizabal, 15San Sebastian 20018Spain

Chapter 9

Dr Yongli LiKey Laboratory of AdvancedFunctional Materials, Ministry ofEducation of ChinaBeijing University of TechnologyPingleyuan 100, Chaoyang DistrictBeijing 100022China

Tel: +86 10 67391760,Fax: +86 10 67392840E-mail: [email protected],[email protected]

Chapter 10Dr Yoshiaki MorisadaJoining and Welding ResearchInstituteOsaka UniversityIbarakiOsaka 567-0047Japan

Tel: +81-6-6879-8693E-mail: [email protected]

Professor Yoshinari Miyamoto*Joining and Welding ResearchInstituteOsaka UniversityIbaraki,Osaka 567-0047Japan


Chapter 11

Professor Wayne D. Kaplan* andDr Amir AvishaiDepartment of MaterialsEngineeringTechnion – Israel Institute ofTechnologyHaifa 32000Israel


Contributor contact details xv

Chapter 12

Professor Alain Peigney*CIRIMAT, UMR UPS-INPT-CNRS5085Université Paul-Sabatier, Bat 2R1118 Route de Narbonne31062 Toulouse Cedex 4France

Tel: 05 61 55 61 75E-mail: [email protected]

Professor Christophe LaurentCIRIMAT, UMR UPS-INPT-CNRS5085Université Paul-Sabatier, Bat 2R1118 Route de Narbonne31062 Toulouse Cedex 4France

Tel: 05 61 55 61 63E-mail: [email protected]

Chapter 13

Dr Ruigang WangScience and Engineering ofMaterials Program and Center forSolid State ScienceArizona State UniversityTempe, AZ 85287-1704USA


Chapter 14

Assistant Professor Dr CemailAksel*Anadolu UniversityFaculty of Engineering andArchitectureDepartment of Materials Scienceand EngineeringIki Eylül Campus26470 EskisehirTurkey

Tel: 00-90-222-3350580/6362Fax: 00-90-222-3239501E-mail: [email protected]

Professor Frank L. RileyDepartment of Materials, School ofProcess, Environmental andMaterials EngineeringUniversity of LeedsLeeds LS2 9JTUK

Tel: 00-44-113-3432531Fax:00-44-113-2422531E-mail: [email protected]

Chapter 15

Dr Christos Kastritseas, Dr PaulSmith and Dr Julie Yeomans*Reader in Ceramic MaterialsSchool of EngineeringPostal Area H6University of SurreyGuildfordSurrey GU2 7XHUK

Tel: +44 (0) 1483 689613E-mail: [email protected]

Contributor contact detailsxvi

Chapter 16

Arturo Domínguez-Rodríguez* andD. Gómez-GarcíaDepartamento de Física de laMateria CondensadaUniversidad de SevillaApartado 106541080 SevillaSpain


F. WakaiMaterials and StructuresLaboratoryTokyo Institite of Technology4259 NagatsutaMidoriYokohama 226-8503Japan

Chapter 17

Professor Servet Turan*Anadolu UniversityDepartment of Materials Scienceand EngineeringIki Eylül Campus26555 EskisehirTurkey


Dr Kevin M. KnowlesUniversity of CambridgeDepartment of Materials Scienceand MetallurgyPembroke StreetCambridge CB2 3QZUK


Chapter 18

Dr Z. B. Yu*Centre for Manufacturing ofAdvanced Ceramics andNanomaterialsQueen’s UniversityKingston,Ontario, K7L 3N6,Canada


D. P. ThompsonAdvanced Materials GroupSchool of Chemical Engineeringand Advanced MaterialsUniversity of NewcastleNewcastle upon Tyne NE1 7RUUK

E-mail:[email protected]

Chapter 19

Dr Csaba BalázsiCeramics and CompositesLaboratoryResearch Institute of TechnicalPhysics and Materials ScienceHungarian Academy of Sciences1121 Budapest XIIKonkoly-Thege u. 29-33Hungary

Tel: +36-1-3922222/3279Fax: +36-1-3922226http://www.mfa.kfki.huE-mail: [email protected]

Contributor contact details xvii

Chapter 20

Dr Yu ZhangDepartment of Biomaterials andBiomimeticsNew York University College ofDentistryRoom 813C345 East 24th StreetNew York, NY 10010USA


Chapter 21

Professor Stuart HampshireMaterials Ireland Research Centre& Materials and Surface ScienceInstituteUniversity of LimerickLimerickIreland


Chapter 22

Guy Anné, Jozef Vleugels andOmer Van der Biest*Department of Metallurgy andMaterials Engineering (MTM)Katholieke University, Leuven,Kasteelpark Arenberg 44B-3001 Heverlee (Leuven)Belgium


Contributor contact detailsxviii

1

By definition, composite materials are formed from two or more materialsthat have quite different properties. The resultant material has a heterogeneousmicrostructure with extraordinary performance that displays a combinationof the best characteristics of the component materials. Composites are widelyused because their overall properties can be engineered through microstructuraldesign to become superior to those of the individual monolithic counterparts.

Nature has provided some of the best-performing composites such asseashell nacre, bones, macadamia nutshells, wood and bamboo. These naturalcomposites have superior mechanical efficiency in strength, hardness andtoughness compared to many man-made composite materials. These biologicalcomposites display graded structures at several levels of hierarchy with lengthscales that range from micro- to nanometres. For instance, seashells havetwo to three orders of lamellar structure whilst bone has seven orders ofhierarchy. Although these materials have complicated hierarchical structures,the most basic level of hierarchy occurs at the nanoscale where nanometre-sized hard inclusions are embedded within the soft protein matrix. In seashellsand bone, this involves the fine dispersion of staggered aragonite bricks andmineralized fibrils, respectively. This nanostructure is expected to play animportant role in the overall mechanical properties of biological materials. A‘bricks and mortar’ lamellar structure exists in nacre where the thickness ofthe aragonite bricks is several hundred nanometres. A well-defined organizationof microstructure exists in seashells in the form of interlaced bricks (mineralplatelets) ‘glued’ together by protein layers. The length scale of the mineralplatelets is normally in the range of 100–500 nm. Following similar principles,the cell walls of wood or bamboo are made of hard cellulose fibrils embeddedin a soft hemicellulose–lignin matrix.

At the level of individual mineralized fibrils, the mechanical properties ofnatural composites depend on the precise arrangement of mineral crystalswithin the fibrils as well as functional loading. It is known that in themicrostructure of bone, the mineral crystals have a large aspect ratio and are

Introduction

I M L O W , Curtin University of Technology, Australia

Ceramic matrix composites2

aligned along collagen fibres with a preferred orientation parallel to thelongitudinal axis of the bone. The size and geometry of the mineral crystalsare thought to play important roles in the mechanical properties of biomaterials.For example, the arrangement of mineral platelets in preferred orientationsproduces intrinsically anisotropic biocomposites. These anisotropic propertiesfacilitate further adaptation to the environmental loading in natural evolution.Bone can possess an anisotropic ratio of 1.7–2.1 in two normal directions,whilst seashells have strongly anisotropic mechanical properties accordingto their lamellar macro- and microstructure. These studies have demonstratedthat mineral crystal arrangement is crucial for the strength and fractureproperties of bone.

Stiffness, strength, toughness and hardness are very important in relationto the supporting and protecting functions of natural composites. Whilst thematrix phase is softer and weaker than the dispersed phase by several ordersof magnitude, the stiffness and strength of resultant composites is notsignificantly reduced through the presence of matrix such as lignin or protein.Furthermore, since the matrix is relatively soft with little capability forsustaining an external load, the dispersed phase must therefore be assumedto carry most of the external load acting on the composite. However, thedispersed phase is very brittle and thus the matrix plays an essential role inachieving and maintaining a high degree of toughness. Qualitatively, thematrix behaves like a soft, surrounding layer which protects the dispersedphase(s) from the peak stresses caused by the external load and homogenizesthe stress distribution within the composite. It is interesting to note that thematrix content has a profound influence on the functionality and mechanicalproperties of natural composites. For instance, the volume fraction of proteinis approximately 50% in bone but only 5% in nacre and 3% in bamboo. Thisdifference may be attributed to the fact that the dynamic loading anddeformation in these materials are function specific.

The use and development of composite materials through the mimickingof nature has been a part of mankind’s technology since the first ancientbuilder used straw to reinforce mud bricks for improved strength and toughness.Later, the Mongols in the twelfth century made the most advanced weaponswith archery bows that were smaller and more powerful than those of theirenemies. These bows were made of composite structures by combining cattletendons, horn, bamboo and silk, bonded with natural pine resin. The tendonswere placed on the tension side of the bow, the bamboo was used as a core,and sheets of horn were laminated to the compression side of the bow. Theentire structure was tightly wrapped with silk using the rosin adhesive toachieve high strength and impressive performance. More recently, concretewith good compressive strength has been fabricated by dispersing aggregatesin cement. Reinforced concretes with stronger tensile strength have beenproduced using metal rods, wires, mesh or cables.

Introduction 3

The advent of engineering-designed polymer matrix composites (PMCs)in the late 1940s has provided an impetus for the emergence of sophisticatedmetal-matrix (MMCs) and ceramic-matrix composites (CMCs). Althoughthere are many similarities between these three classes of composites, especiallyin terms of the physics of strengthening and toughening, the chemical processesinvolved in the interfacial reactions and fabrication of CMCs (and MMCs)are much more complex and challenging. Again, nature has provided thefundamental concept or inspiration for the development of these unique andhigh-performance synthetic materials. In particular, the study of naturalcomposites has provided material scientists and engineers with a bettermechanistic understanding for the design of new composite materials withsuperior mechanical efficiency in strength, toughness and stiffness. Indeed,super-tough ceramic composites with a laminated microstructure have beendesigned through the bio-mimicry of nacre shells, bamboo and nutshells.

The manipulation of microstructures at the nano- or micro-level is the keyto controlling the properties of all materials, including ceramic matrixcomposites. The profound interdependence of microstructure and propertiesis central to all forms of composites irrespective of the matrices. Thisdependency allows the microstructure of a composite to be tailored at theprocessing stage to achieve the desired physical, chemical or mechanicalproperty for a specific application.

It is widely recognized that the properties of CMCs are controlled by thesize and volume fraction of the reinforcements as well as the nature of thematrix–reinforcement interfaces. An optimum set of mechanical propertiescan be achieved when fine particulates or fibres are dispersed uniformly inthe matrix. This has led to the development of in-situ reinforced compositesin which the reinforcements are synthesized in the matrix by chemical reactionsbetween elements or between element and compound during the compositefabrication. Compared to the conventional composites produced by ex-situmethods, the in-situ or self-reinforced composites exhibit superior properties,which include less degradation in elevated-temperature services by virtue ofenhanced interfacial stability, stronger interfacial bonding and better mechanicalproperties because of finer and more uniform dispersion of the reinforcements.

The mechanical performance of CMCs is also sensitive to the phasecomposition and morphology of their microstructures. For instance, ahomogeneous and fine-grained alumina offers strength, wear and fatigueresistance, but is brittle and not damage tolerant. By judiciously adjustingthe composition by addition of an appropriate amount of titanium oxide, aduplex microstructure of uniformly dispersed aluminium titanate, AT (Al2TiO5),can be produced with improved flaw tolerance. Further improvement in flawtolerance can be achieved through uniform incorporation of large aluminagrains within a fine-grained alumina–AT matrix (duplex-bimodal


microstructure). Although such heterogeneous microstructures can impartconsiderable increases in damage resistance, they degrade the intrinsic strengthand wear resistance.

A different but more effective approach to achieve flaw tolerance withoutsubstantial strength degradation is based on in-situ growth of aluminateplatelets (e.g. LaAl11O18, LaMgAl11O19, SrAl12O19, CaAl12O19 andNa2MgAl12O17) within the alumina matrix. Strengths of over 600 MPa andfracture toughness in excess of 6 MPa.m1/2 have been achieved in aluminacomposites with both elongated Al2O3 grains and LaAl11O18 platelets.

Recent developments in layered ceramics have also provided a strategyfor designing laminates with high strength and high toughness. These structureshave an outermost homogeneous layer to provide strength or wear resistance,and an underlying heterogeneous layer to provide toughness. Unlike moretraditional layered structures which promote toughness by interlayer crackdeflection or strength by incorporating macroscopic compressive residualstresses, the new approach deliberately seeks to produce strong interlayerbonding and to eliminate residual macroscopic stresses. Accordingly, anyattendant counterproductive effects of weak interlayers and residual stressesfrom delamination failures can be avoided. Layered ceramics produced inthis way at the National Institute of Standards and Technology have shownuncommonly high damage resistance under Hertzian loading, with retentionof both strength and wear resistance.

Similar advances with outstanding mechanical performance have alsobeen made through the microstructural modification of ceramic matrices viathe dispersion of whiskers, fibres and nano-sized particles. In brittle ceramicmatrix composites with fibre reinforcements, the favourable fracture behaviouris provided by the presence of weak fibre–matrix interfaces, which lead tothe well-known fibre pullout effect.

The main advantage of CMCs over their monolithic counterparts lies inthe fact that they are tough although their constituents are intrinsically brittle.This key property is achieved through a proper design of the fibre–matrixinterface arresting and deflecting cracks formed under load in the brittlematrix and preventing the early failure of the fibrous reinforcement. From amechanical standpoint, these CMCs are damageable elastic materials, i.e.when loaded at a high enough level, microcracking and interfacial debondingoccur, which are responsible for reduced stiffness and non-linear stress–strain behaviour. On the one hand, these damaging phenomena are beneficialsince they are at the origin of the non-brittle character of these ceramics. Onthe other hand, they are detrimental for non-oxide CMCs since they favourthe in-depth diffusion of oxygen towards the oxidation-prone interphase andfibres which in turn may embrittle the composites.

Hybrid CMCs have also been created by incorporating simultaneouslyfibres, particles and/or whiskers as reinforcing elements. In hybrid CMCs,

Introduction 5

the fibres are used to impart fracture resistance and toughness by exploitingwell-known fracture mechanisms such as fibre pullout and crack deflection.The particulate phase is used to improve other engineering properties relevantfor the intended application of the materials, such as thermal shock resistance,wear resistance or impact strength.

The development of CMCs is a promising means of achieving lightweight,structural materials combining high-temperature strength with improvedfracture toughness, damage tolerance and thermal shock resistance.Considerable research effort is being expended in the optimization of ceramiccomposite systems, with particular emphasis being placed on the establishmentof reliable and cost-effective fabrication procedures.

This book consists of a collection of chapters reviewing and describingthe latest advances, challenges and future trends in the microstructure andproperty relationship of various CMCs. Each chapter has been written byexperienced and established researchers who have worked for many years intheir respective fields. The book covers all aspects of the interdependence ofprocessing, microstructure, properties and performance of the following fivecategories of CMCs:

∑ Fibre-whisker- or particulate-reinforced CMCs∑ Graded and layered CMCs∑ Nanostructured CMCs∑ Non-oxide CMCs∑ Refractory and speciality CMCs.

Bibliography

An, L., Chan, H.M., Padture, N.P. and Lawn, B.R. Damage-resistant alumina-based layercomposites. J. Mater. Res. 11 (1996) 204.

Chen, P.L. and Chen, I.W. In-situ alumina/aluminate platelet composites. J. Am. Ceram.Soc. 75 (1992) 2610.

Feng, Q.L., Cui, F.Z., Pu, G., Wang, R.Z. and Li, H.D. Crystal orientation, tougheningmechanisms and a mimic of nacre. Mater. Sci. Eng. C 11 (2000) 19–25.

Harmer, M., Chan, H.M. and Miller, G.A. Unique opportunities for microstructuralengineering with duplex and laminar ceramic composites. J. Am. Ceram. Soc. 75(1992) 1715.

Hou, D.F., Zhou, G.S. and Zheng, M. Conch shell structure and its effect on mechanicalbehaviour. Biomaterial 25 (2004) 751–756.http://www.acmanet.org/professionals/index.cfm

Ji, B. and Gao, H. Mechanical properties of nanostructure of biological materials. J.Mech. Phys. Solids 52 (2004) 1963.

Katti, D.R., Katti, K.S., Sopp, J.M. and Sarikaya, M. 3D finite element modeling ofmechanical response in nacre-based hybrid nanocomposites. Comput. Theoret. Polym.Sci. 11 (2001) 397–404.

Li, S.H., Liu, Q., de Wijn, J.R., Zhou, B.L. and de Groot, K. In vitro calcium phosphateformation on a natural composite material, bamboo. Biomaterial 18 (1997) 389–395.


Padture, N.P., Bennison, S.J. and Chan, H.M. Flaw-tolerance and crack resistance propertiesof alumina–aluminium titanate composites with tailored microstructures. J. Am. Ceram.Soc. 76 (1993) 2312.

Yasuoka, M., Hirao, K., Brito, M.E. and Kanzaki, S. High strength and high toughnessceramics in the Al2O3/LaAl11O18 systems. J. Am. Ceram. Soc. 78 (1995) 1853.

Part I

Fibre-whisker- and particulate-reinforcedceramic composites

7

9

1.1 Introduction

Fibrous monolithic ceramics (FMs) consist of a hexagonal arrangement ofsubmillimeter ‘cells’ of strong polycrystalline ceramic and a network ofcrack-deflecting weak ‘cell boundaries’ [1, 2]. These composites are sinteredor hot-pressed monolithic ceramics with a distinct fibrous texture. This uniquearchitecture opened new avenues for ceramic composites, in which they failin a nonbrittle manner because of crack interactions with weak cell boundaries,such as crack deflection or crack delamination. This approach provides asimple and versatile method for manufacturing nonbrittle ceramic compositesfrom a variety of different material combinations that include oxide ceramics(Al2O3/Al2O3–ZrO2 [3]) and non-oxide ceramics (SiC/graphite [4, 5], SiC/BN [6] and Si3N4/BN [1, 7–16]).

This chapter presents an overview of these composites with a variety ofmaterial combinations, as well as their architectures. The major objectives ofthis chapter are the following: (a) to briefly discuss the history and mainconcepts of FMs in section 1.2; (b) to address the experimental ways toprepare these composites, including coextrusion, microfabrication bycoextrusion, and hybrid extrusion and dip-coating, in section 1.3; (c) todemonstrate the various types of FMs in the form of oxide-based or non-oxide-based composites, as well as their various architectures in section 1.4;(d) to discuss the mechanical properties of Si3N4/BN FMs at room and hightemperatures, as well as their fracture mechanisms, in section 1.5; and finally(e) to give a personal perspective on the future of these wonderful compositesin section 1.6.

1.2 History

Ceramics have been well known to offer potential benefits over metal partsin high-temperature environments, such as higher strength, lower density,and greater resistance to oxidation. In spite of these merits, their low fracture

1Fibrous monolithic ceramics

Y - H K O H, Seoul National University, Korea


resistance with poor reliability has long been a major concern. A number ofpotential toughening mechanisms have been proposed [17, 18]. One of themost successful techniques is the fiber-reinforced ceramic composites [19,20]. In a tough ceramic composite, the matrix crack is deflected at theinterface and ceases to propagate, allowing the composite to possess hightoughness along with nonbrittle failure by extensive crack interactions, suchas bridging of the primary crack, crack deflection along the fiber/matrixinterface and frictional sliding with fiber pullout of the matrix [20]. However,fiber-reinforced ceramic composites have not been fully utilized yet becauseof their manufacturability and cost. Also, densification techniques such as(hot pressing and hot isostatic pressing) often lead to fiber damage, deterioratingits mechanical properties.

Another toughening concept was introduced by Cook and Gordon in 1964[21 They suggested crack propagation in brittle materials could be controlledby incorporating a fabric of microstructural features that change the crackpath, resulting in high toughness by crack blunting at the weak interface andcrack delamination. Rather than rely on bridging mechanisms to improvetoughness, the advancing crack was blunted at the weak interface and forcedto reinitiate in order to continue propagation. Based on this concept, in 1988Coblenz described a method for producing a pressureless sinterable ceramiccomposite consisting of a strong load-bearing phase surrounded by a continuousweak crack-deflecting phase [2].

Another variation on the concept of Cook and Gordon was introduced byClegg et al. in 1990 [22]. This approach involves alternating layers of thestrong load-bearing phase and the weak crack-deflecting phase, allowing thecomposite to fail in a nonbrittle manner with high strength and toughness(Fig. 1.1(a)). Following failure of this layer, a series of decreasing steps wereobserved in the load–deflection behavior of the laminate, as additional layersfailed with subsequent crack arrest via deflection along the weak interfaces.

In 1991, Halloran and co-workers pioneered an exciting new class ofstructural ceramics, so-called ‘fibrous monolithic ceramics’ that exhibitmechanical properties similar to ‘continuous fiber ceramic composites’,including very high fracture energies, damage tolerance and graceful failure[1, 3–5]. They consist of 250-micron ‘cells’ of a strong polycrystalline ceramic,such as silicon carbide or silicon nitride, separated by ‘cell boundaries’ frommaterials, such as boron nitride, which promote crack deflection anddelamination (Fig. 1.1(b)). Fibrous monolithic ceramics are produced mostoften by extrusion, followed by lay-up of filaments into laminates. Theextruded filaments consist of a cell phase surrounded by a sheath that formsa continuous cell boundary. This approach creates analogs of many compositearchitectures, allowing these composites to fail in a nonbrittle manner withenergy dissipation arising from sliding of the cells, and branching and deflectionof cracks. Such fibrous monolithic ceramics constitute lower-cost alternatives

Fibrous monolithic ceramics 11

to conventional continuous-fiber ceramic composites in some applications,and a wide variety of fibrous monolithic ceramics are available commercially.A number of the different cell/cell-boundary combinations have beeninvestigated, including oxide and non-oxide ceramics.

1.3 Processing

In this section, we discuss three prevalent processing methods for producingfibrous monolithic ceramics, i.e. coextrusion [1, 23], microfabrication bycoextrusion [24], and hybrid extrusion and dip-coating [25].

1.3.1 Coextrusion

A significant advancement in the processing of FMs was made by thedevelopment of coextrusion. The schematic illustrations showing processingare shown in Fig. 1.2. This process involves forming a feedrod consisting ofa core of the cell material surrounded by a shell of the cell boundary material,prepared using a cylindrical mold and a half-pipe shape (Fig. 1.2(a)). Boththe core and the shell are blends of thermoplastic polymer and ceramicpowder. The feedrod is then used to coextrude a ceramic green fiber identicalto the feedrod in core/shell proportion, but 100 times smaller (Fig. 1.2(b)).The coextruded fibers are uniaxially aligned and then warm pressed to fabricategreen billet (Fig. 1.2(c)). In addition, for the multilayer FMs, the filamentdirection can be rotated between the layers. The resulting green billets undergobinder removal and hot-pressing to produce densified FMs (Fig. 1.2(d)).

(a) Laminated composite

(b) Fibrous monolithic ceramic

Strong layer

Weak layer

Strong cell

Weak cell boundary

1.1 Schematics illustrating (a) laminated composite, consisting of thestrong layer and the weak layer, and (b) fibrous monolithic ceramic,consisting of the strong cell and the weak cell boundary.


There are many advantages to this technique. In general, the coatings onfibers produced by this method are more uniform than those on dip-coatedfibers, improving the overall uniformity of the composite. Unlike dip-coating,the coating thickness is determined by the shell thickness on the feedrod, notthe rheological properties of the dip-coating slurry. Thus batch-to-batchrepeatability is also improved. The coextruded fibers are also much easier tohandle than the dry-spun green fibers, making processing much easier toperform.

1.3.2 Microfabrication by coextrusion (MFCX)

A variation on this approach used multifilament coextrusion, so-called‘microfabrication by coextrusion (MFCX )’. A limitation of the single-filamentprocess is the size of the filament. The rheological properties of the polymer/ceramic blends make spinning fibers smaller than 250 mm very difficult.Additionally, spooling fine-diameter fibers is quite challenging. The MFCXis shown schematically in Fig. 1.3. The setup is the same as that used to spinfibers except that the spinneret is replaced with an extrusion die with adiameter between 1 mm and 6 mm. Two separate extrusion steps are used. Inthe first step, coarse primary filaments are extruded from the feedrod (Fig.

Cell boundary

(a)

(b)

(d)(c)

Hot-pressing

Cell

1.2 Schematic illustrations showing coextrusion process to fabricateFMs: (a) initial feedrod formation, (b) coextrusion, (c) lamination, and(d) hot-pressing.


1.3(a)). The primary filaments are bundled together to form a second feedrod.On extrusion of the second feedrod, a filament containing many cells in itscross-section is formed (Fig. 1.3(b)). The second filament is referred to as amultifilament strand, indicating the origin of the strand (Fig. 1.3(c)). Thesimilar procedure to coextrusion is employed to fabricate specimens, includingstrand alignment, lamination, and thermal treatment.

The scale of the cells within the multifilament strand is controlled by twofactors: (1) the number of primary filaments bundled into the second feedrod,and (2) the size of the extrusion die used to form the second filament. Thenumber of filaments bundled into the second feedrod is determined by boththe size of the primary filaments and the size of the second feedrod. Thedecreasing ratio of extrusion die sizes results in a dramatic increase in thenumber of cells within a strand. In addition to making finer cell sizes possible,it is also much easier to lay coarse multifilament strands into a die than it isto lay small fibers. The strands lay straight and flat, whereas fibers tend tocurl up and become intertwined. However, despite the fine cell size, thecoarseness of the strands limits their use to architectures in which the scaleof a cluster of cells is no less than ~ 0.75 mm.

1.3.3 Hybrid extrusion and dip-coating

As previously mentioned, FMs were first fabricated by hybrid extrusion anddip-coating [4–6]. Firstly, highly concentrated, viscous slurries are extrudedand then dried to form a green fiber. The resulting green fibers are dip-coated

(a)

Cell Cellboundary

Primary filament

(b)

Second filament

(c)

1.3 Schematic illustrations showing MFCX process to fabricate FMs:(a) coextrusion, (b) assembly of primary filament, and (c) filamentafter second coextrusion.


in a slurry containing the cell boundary material. The coated fibers (core/shell) are cut to length and stacked in a die to form a green body. The greenbody undergoes the binder removal and hot-pressing.

Although this process limits the uniform coating of cell boundaries, it cancreate FMs containing gradient cell boundary thickness in a simple way. Forinstance, in 2002, Koh et al. fabricated the three-layered Si3N4/BN FMs bysimply adjusting the BN-containing slurry for the cell boundary phase duringthe dip-coating stage (Fig. 1.4) [25]. In this process, Si3N4-polymer doughinstead of thermoplastic ceramic compound is used for spinning. For the cellboundary, two kinds of dip-coating slurries with different BN concentrationswere used. The Si3N4-polymer is extruded using a syringe with a pistonthrough a 300 mm orifice, and then dip-coated by passing the BN-containingslurry with 20 wt% and 0 wt% concentration (Fig. 1.4(a)). The coated fiberwas uniaxially arrayed and then dried in an oven at 80∞C for 12 h to improvethe shape and strength of the fiber by hardening. The green billet was pressedusing a square mold. Thereafter, the green billet undergoes thermal treatmentto produce FM composite, consisting of monolithic Si3N4 and Si3N4/BN FM(Fig. 1.4(b)).

1.4 Structures

1.4.1 Various material combinations

Fibrous monolithic ceramics consist of dense cells separated by a continuouscell boundary, in which the cells provide most of the strength of the FM and

Monolithic Si3N4

Si3N4/BN FM

Monolithic Si3N4

(b)

Si3N4polymer

0 wt% 20 wt%

BN-containing slurry

(a)

1.4 (a) Schematic illustrating hybrid extrusion and dip-coatingprocess to fabricate a Si3N4/BN FM composite; (b) schematic of theFM composite, consisting of monolithic Si3N4 and Si3N4/BN FM.


the cell boundary provides the toughness by isolating the cells from eachother and promoting dissipation of fracture energy by mechanisms such aspullout of the cells or deflection of a crack through the cell boundary [1].The cell boundaries must be either weak themselves or poorly bonded to thecells to dissipate fracture energy and exhibit minimal or no reaction with thecells for long-term use at elevated temperatures. To date, many kinds ofstructural ceramics have been examined for a strong cell phase. They are inthe forms of either oxides (Al2O3 [3], ZrO2 [26], and ZrSiO4 [27]) or non-oxides (SiC [4–6], Si3N4 [8], and borides [1, 7–16]). Oxides have the advantageof stability in oxidizing environments, while non-oxides have the advantageof substantially higher strength and superior creep resistance.

In all-oxide FMs, porous cell boundaries are generally employed becausethey can minimize transmission of fracture energy to the cells. They havebeen formed from large-grained ceramic powders or a mixture of ceramicparticles and platelets. In addition, the cell-boundary phase must be resistantto microstructural alteration with prolonged heating for elevated-temperatureapplication. For non-oxide FMs, similar criteria for selecting a cell boundaryshould be considered. To date, hexagonal-BN has been extensively used ascell boundary because BN forms a dense, highly textured cell boundary thatbonds only weakly to the cell by hot-pressing. In most cases, Si3N4 has beenused for the cells. Oxide sintering aids (e.g. Al2O3 and Y2O3) in the Si3N4

migrate into the BN cell boundary during hot-pressing and impart the bonding[1, 24].

1.4.2 Various architectures

The processing methods used to manufacture FMs are very versatile, allowingany number of unique, submillimeter architectures to be designed andfabricated. For example, sheets of filament can be rotated with respect to oneanother to form a multiaxial architecture (Fig. 1.5(a)) [1, 25]. Another goodexample is to tailor the cell boundaries, in which the cell boundaries containa thin web of Si3N4 reinforcement (Fig. 1.5(b)) [1, 25]. Such Si3N4

reinforcement may alter the fracture behaviors and mechanical properties ofSi3N4/BN FMs at both room and high temperatures. This capability of tailoringthe architecture can offer the opportunity to allow FMs to achieve betterperformance at both room and high temperatures.

1.5 Mechanical properties

Since fibrous monolithic ceramics are intended for use in applications wherestresses are primarily generated due to bending, strength and work-of-fracturein flexure are measured to evaluate their basic mechanical properties. Inaddition, factors determining the manner of crack propagation should be


examined, including the fracture resistance of cells, the interfacial fractureresistance, and the interfacial sliding resistance, and residual stresses shouldbe investigated to understand how cracks propagate in flexure. Among themany FMs that have been investigated, Si3N4/BN FMs have achieved thebest overall properties and have been reliably manufactured on a commercialscale [28]. Therefore, this section deals with the mechanical properties ofSi3N4/BN FMs primarily at room and high temperatures, as well as withtheir failure mechanisms.

1.5.1 Room-temperature properties

Mechanical properties and failure mechanisms

Si3N4/BN FMs have shown promise as a structural and tough ceramic material[1, 7–15, 29]. The unique flexural response of FMs is a result of thesubmillimeter architecture that is engineered into the material. A typicalmicrostructure of fibrous monolithic Si3N4/BN ceramic (FM) is shown inFig. 1.6. Low-magnification SEM micrographs of polished sections showthree-dimensional representations of the submillimeter structure of twoarchitectures of fibrous monoliths. The polycrystalline silicon nitride cells(~250 mm) appear in dark contrast, while the continuous boron nitride cellboundaries (20 mm) appear in bright contrast. The Si3N4 cells are surroundedby the cell boundaries consisting of BN particles bonded with yttriumaluminosilicate.

The typical stress versus crosshead deflection response associated with aSi3N4/BN FM is contrasted with that of a monolithic Si3N4 in Fig. 1.7. Themonolithic Si3N4 exhibits greater strength but negligible apparent work-of-fracture (WOF) because of brittle failure (Fig. 1.7(a)). On the other hand, theSi3N4/BN FM demonstrates some load-bearing ability following the first

(b)

(a)

50 mm

Si3N4

BN

Si3N4 Cell

1.5 (a) Low-magnification SEM composite showing three sections ofa fibrous monolith with a [0∞/90∞] architecture; (b) cross-sectionalview of FM fabricated with a thin web of Si3N4 reinforcing the BNcell boundary (adapted from ref. [1]).


failure event with significant WOF (Fig. 1.7(b)). To date, the flexural strengthand WOF of Si3N4/BN FMs have significantly improved up to ~700 MPaand ~8000 kJ/m2, by optimizing respectively manufacturing routes andmicrostructures [1].

In unidirectional FMs, fracture can initiate when the tensile stress carriedby the cell exceeds the strength of the cell [1]. This is generally favorablewhen the cell boundaries are tough in comparison to those of the cells. A

250 mm

1.6 Low-magnification SEM micrograph of polished sections,showing three-dimensional representations of the submillimeterstructure of two architectures of Si3N4/BN FM (adapted from ref. [29]).

0.0 0.2 0.4 0.6 0.8 1.0Crosshead displacement (mm)

800

600

400

200

0

Ap

par

ent

stre

ss (

MP

a)

(a) Monolithic Si3N4

(b) Fibrous monolith

1.7 Flexural response of (a) monolithic Si3N4 and (b) Si3N4/BN FM(adapted from ref. [29]).


flaw on the tensile surface of FMs will cause failure of only a single cell. Themaximum applied load is typically achieved at the point just prior to failurein the layer of the cells closest to the tensile surface, because the load-bearing capacity is reduced due to the reduction in the effective cross-sectionof the sample. Thereafter, each stress drop in FMs is frequently followed,owing to the fracture of one or several layers of cells (Fig. 1.7(b)).

As well as the strength of the cell, the strength of unidirectional FMsdepends on their orientation with respect to the loading axis. The flexuralstrength of FMs decreases dramatically as the direction of stress applicationchanges from on-axis to off-axis [1, 24]. The strength of FMs tested off-axisis much lower than that for on-axis orientation because failure is determinedby the strength of the cell boundaries, rather than that of the cells. FMs withmultiaxial architectures will be an important part of designing for complexstress states because many applications involve biaxial or off-axis loadingconditions. The strength is determined by both cell and cell boundary fractures.If cells on the tensile surface are not aligned in the direction of the appliedstress, failure of the cell boundary on the tensile surface can occur at arelatively low load, but cells with 0∞ orientations that are just beneath thetensile surface can continue to bear substantially more load [1, 24].

In addition to strength and WOF of FMs, the elastic behavior of thesearchitectures should be considered. Simple brick models were proposed toaccurately predict elastic properties of FMs [1, 24]. Figure 1.8 shows theelastic modulus versus orientation for uniaxially aligned Si3N4/BN FMswith experimentally measured values, indicating that there is very goodagreement between experiment and prediction. This prediction can be usedfor FMs with multiaxial architectures.

MeasuredPredicted

0 20 40 60 80Orientation, q (degrees)

300

250

200

150

100

You

ng

’s m

od

ulu

s E

1 (G

Pa)

1.8 Young’s modulus versus orientation for uniaxially aligned Si3N4/BN FM (adapted from ref. [1]). The line is the predicted behaviorusing the brick model and laminate theory.


Crack propagation and energy absorption

The unique nonbrittle failures of FMs are accomplished as a consequence ofcrack deflection at the BN cell boundaries as well as significant crackdelamination and sliding, as shown in Fig. 1.9. During crack propagation,tensile-initiated cracks are turned on weak interfaces, creating surface areaand thus absorbing energy in the process. This process, known as crackdeflection and crack delamination, is repeated as the crack works its waythrough the thickness of the sample.

There are several factors that determine the degree of the crack interactionswith the BN cell boundaries. In general, crack deflection is predicted whenthe fracture resistance of the interface is low and when the elastic mismatchbetween the cell and the cell boundary is high. In 1998, Kovar and co-workers. observed that the addition of a Si3N4 phase into a BN interphase inlaminate composite decreased the interfacial fracture resistance, reducingthe degree of crack delamination [30], as shown in Fig. 1.10(a)–(d). Extensivedelamination and high energy absorption are observed only in materials thathave the lowest interfacial fracture resistance (Fig. 1.10(a)). In addition toSi3N4/BN laminate composites, the fracture resistances of Si3N4/BN FMswere also examined [31, 32].

The energy-absorption capacity of FMs is primarily influenced by thecrack path after the initial crack deflection occurs. Long delamination distancesare favored when the interfacial fracture resistance is low, the flaw size inthe layers is small, and the fracture resistance of the layers is high [30, 33].A map of crack propagation behavior in Si3N4/BN FM is shown in Fig. 1.11.At very low values of the interfacial fracture resistance, increasing the interfacialfracture resistance causes more energy to be dissipated through the creation

1 mm

1.9 Optical photograph of crack propagation of the fibrous monolithicSi3N4/BN ceramic after flexural testing (adapted from ref. [29]).


of interfacial crack area. However, if the interfacial fracture resistance is toohigh, crack kinking will reduce the delamination area. This observation suggeststhat there is an optimum interfacial fracture resistance that maximizes theenergy-absorption capability, and this optimum value is determined by thetransition from delamination and crack kinking.

(a)

(b)

(c)

(d)

1.10 SEM micrographs of the side surface of flexural specimenscontaining (a) 10, (b) 25, (c) 50, and (d) 80 vol% Si3N4 in theinterphase (adapted from ref. [30]).

Crack kinking

Delamination

0 0.1 0.2 0.3 0.4 0.5 0.6

200

150

100

50

0

Cri

tica

l fl

aw s

ize,

a (

mm)

G Gi Si3N4/

1.11 Critical flaw size predicted to cause a crack to kink out of the BNinterphase, plotted against the ratio of the interfacial fractureresistance to the fracture resistance of Si3N4 (adapted from ref. [30]).


1.5.2 High-temperature properties

Elevated temperature mechanical properties

Since Si3N4/BN FMs are candidates for high-temperature applications, forinstance in gas turbine engines, the failure behavior of FMs at various elevatedtemperatures must be understood. In 2000, Trice and co-workers measuredthe high-temperature flexural strength of Si3N4/BN FMs and monolithic Si3N4

from room temperature through 1400∞C [11]. A larger drop in flexural strengthoccurs in the monolithic sample from room temperature to 1000∞C as comparedto the FM sample, as shown in Fig. 1.12. The difference in strength trendsthrough 1000∞C between the monolithic and FM samples may be attributedto the different amounts of glassy phase in the Si3N4 grain of two materials.

They also observed three different failure modes depending on testingtemperatures, as shown in Fig. 1.13. As previously mentioned, the apparentstress peak is related to tensile crack initiation in the Si3N4 cell and thesubsequent drop is related to crack deflection and delamination in the BNcell boundary. In relatively low-temperature regimes (25∞C to 1000∞C), tensile-initiated failure occurred on the outer ply of the Si3N4 cells. However, a sideview of the flexure sample at 1000∞C showed more crack delamination thanat 25∞C, as shown in Fig. 1.14(a) and (b). This change is probably a result ofthe changing interfacial fracture energy [10, 11]. In middle-temperature regimes(1100∞C to 1300∞C), shear-initiated failure occurred at the beam midplane inthe BN cell boundary. Following the peak strength and subsequent reloadingof the specimen, a distinct slope was observed in the apparent stress versuscrosshead deflection curve. The presence of a lengthwise horizontal crackthat split part of the beam into halves, characteristic of midplane-initiated

6Y/2AI-M

6Y/2AI-FM

0 200 400 600 800 1000 1200 1400Temperature (∞C)

1000

800

600

400

200

0

Pea

k st

ren

gth

(M

Pa)

1.12 Comparison of flexural strengths of Si3N4/BN FM and monolithicspecimens with 6 wt% Y2O3 and 2 wt% Al2O3 sintering additives(adapted from ref. [11]).


shear failure, is visible in Fig. 1.14(c). At relatively high temperatures (1400∞C),viscoelastic flow of the Si3N4 on the tensile surface took place. The test barindicated an appreciable bending moment, presumably due to the easy flowof the grain boundary glassy phase present within the Si3N4 cells, as shownin Fig. 1.14(d).

Oxidation behavior

Another critical factor affecting mechanical properties of Si3N4/BN FMs atelevated temperatures is their oxidation resistance. In 2002, Koh and co-workers studied the oxidation behavior of Si3N4/BN FMs after exposure toair at temperatures ranging from 1000∞C to 1400∞C for up to 20 h [34]. Theyobserved an overall weight loss due to the extensive vaporization of B2O3

liquid that was formed by oxidation of BN cell boundary. During oxidation,the BN cell boundary began to oxidize at temperatures above 1000oC, formingB2O3 liquid, followed by its vaporization as oxidation proceeded (not shown),and the Si3N4 cell oxidized at above 1200∞C, leaving Y2Si2O7, SiO2

(cristobalite) crystals surrounded by a glassy phase (Fig. 1.15(a)). At 1400∞C,the Si3N4 cells were also severely damaged, revealing an oxide layer composedof large Y2Si2O7 crystals and a glassy phase, as shown in Fig. 1.15(b). Thesurface cracks were due to crystallization of a glassy phase on cooling.However, the BN cell boundary layers were completely covered by oxidationproducts, implying that the oxidation of Si3N4/BN FM is retarded with furtheroxidation.

Koh and colleagues observed that the maximum apparent strength andWOF decreased with increasing exposure temperature (Fig. 1.16(a) and (b)).

0 1 2 3 4Crosshead deflection (mm)

600

500

400

300

200

100

0

Ap

par

ent

stre

ss (

MP

a)

1400∞C

1300∞C1200∞C

1100∞C

Individualcellbreakage

1000∞C800∞C600∞C

25∞C

1.13 Typical characteristic apparent stress versus crossheaddeflection curves as a function of temperature for Si3N4/BN FM(adapted from ref. [11]).


(a)

(b)

(c)

(d)

1.14 Side view of tested flexure Si3N4/BN FM samples at (a) 25∞C, (b)1000∞C, (c) 1200∞C, and (d) 1400∞C (adapted from ref. [11]).

After severe exposure to air at 1400∞C for 20 h, the sample maintained 41%(~226 MPa) and 21% (~2.3 kJ/m2) of its initial strength and WOF, respectively.These reductions were apparently due to degradation of both the Si3N4 celland the BN cell boundary. In other words, the reduction in strength at


temperatures above 1000∞C was presumably due to the damage of Si3N4

cells on the surface by the tensile-initiated failure, while the reduction inWOF is primarily influenced by the damage of the BN cell boundary withincreasing oxidizing temperatures.

Thermal shock resistance

Some kind of thermal shock loading is inevitable during service of FMs. Inaddition, most FMs have anisotropic thermal-expansion coefficients due totheir unique architectures. In 2004, Koh and co-workers investigated thethermal shock resistance of Si3N4/BN FMs [29]. They observed their excellentthermal shock resistance by measuring the retention of the flexural strengthafter thermal shock test, as shown in Fig. 1.17. The monolithic Si3N4 exhibited

(a)

(b)

30 mm

30 mm

1.15 SEM micrographs of the Si3N4/BN FM samples exposed to airfor 20 h at (a) 1200∞C and (b) 1400∞C (adapted from ref. [34]).


0 200 400 600 800 1000 1200 1400 1600Exposure temperature (∞C)

(b) Work-of-fracture

(a) Flexural strength

800

600

400

200

0

Flex

ura

l st

ren

gth

(M

Pa)

12

10

8

6

4

2

0

Wo

rk-of-fractu

re (kJ/m2)

1.16 (a) Flexural strength and (b) work-of-fracture of the Si3N4/BN FMsample as a function of the exposure temperature (adapted from ref.[34]).

1.17 Flexural response of (a) monolithic Si3N4 and (b) Si3N4/BN FMafter thermal shock test. Monolithic Si3N4 showed brittle fracture,while fibrous monolith showed graceful fracture due to its uniquearchitecture (adapted from ref. [29]).

(a) Monolithic Si3N4

(b) Fibrous monolith

0 200 400 600 800 1000 1200 1400Temperature difference (∞C)

1000

800

600

400

200

0

Flex

ura

l st

ren

gth

(M

Pa)


the typical thermal shock behavior of brittle ceramics, with a critical temperature(DTc = 1000∞C) at which the strength decreases catastrophically [35–38], asshown in Fig. 1.17(a). On the other hand, the flexural strength of Si3N4/BNFM after thermal shock was not changed much without DTc, as shown in Fig.1.17(b), indicating the material’s excellent thermal shock resistance.

They observed similar failure modes for FM samples (i.e., shear-initiatedcracks), regardless of temperature difference (DT), as shown in Fig. 1.18 (a)–(d). With increasing DT, more extensive crack interactions were observed,increasing the apparent WOF (Fig. 1.19). The increase in WOF after thermalshock suggests that thermal shock reduces the interfacial crack resistance ofthe cell boundary, which is a composite of boron nitride and glass. Thermalshock damage seems to be absorbed by pre-existing microcracks within theBN cell boundaries [1, 24], which would decrease the cell boundary fractureresistance, enabling easier crack deflection and higher WOF.

Koh and colleagues suggested models to explain the thermal shock resistanceof Si3N4/BN FMs, as shown in Fig. 1.20(a) and (b). After thermal shock, twothermal stresses are generated in the transverse and longitudinal directionsof the sample, due to its unidirectional architecture (Fig. 1.20(a)). Less thermalstress was generated than in monolithic Si3N4. They also found that thecalculated crack propagation parameter value of the FM sample is muchlarger (>16 times) than that of monolithic Si3N4, indicating that crackpropagation is more restricted for the Si3N4/BN FM sample.

(a) DT = 800∞C (b) DT = 1000∞C

(c) DT = 1200∞C (d) DT = 1400∞C

0.0 0.5 1.0 1.5 2.0 0.0 0.5 1.0 1.5 2.0Crosshead displacement (mm)

300

200

100

0300

200

100

0

Ap

par

ent

stre

ss (

MP

a)

1.18 Flexural strength of Si3N4/BN FM ceramic after thermal shockwith temperature difference DT (adapted from ref. [29]).


0 200 400 600 800 1000 1200 1400Temperature difference (∞C)

14

12

10

8

6

4

Ap

par

ent

wo

rk-o

f-fr

actu

re (

kJ/m

2 )

1.19 Work-of-fracture (WOF) of Si3N4/BN FM ceramic after thermalshock with temperature difference DT (adapted from ref. [29]).

1.20 (a) Thermal stress after thermal shock in transverse andlongitudinal directions; (b) a schematic of single Si3N4 cell betweentwo BN-rich cell boundaries, illustrating (i) pre-existing cracks in cellboundaries (- - -) and cell boundary cracks extended by thermalshock (—), (ii) possible transverse cracks in Si3N4 cell (adapted fromref. [29]).

(a) Thermal stress after thermal shock

slongstrans

(b) Cracking due to thermal stress

100 mm

BN-richcell boundary

BN-richcell boundary

Si3N4 cell

Cell crackingslong ª 0.80sm

strans ª 0.55sm


In addition to these factors, thermal shock-induced cracks and crackpropagation during subsequent flexural testing are also important to explainthe thermal shock resistance of Si3N4/BN FM. The longitudinal thermalstress may fracture occasional Si3N4 cells, and the transverse thermal stressmost likely causes localized extension of the pre-existing flaws in the BN-rich cell boundary and is unlikely to cause cracks within the Si3N4 cells, asshown in Fig. 1.20(b). The extension of BN-cell boundary cracks is believedto decrease the cell boundary fracture resistance, which is consistent with theobservation of more extensive delamination after flexural testing of severelyshocked samples. To clarify this hypothesis, Koh and colleagues measuredthe degree and length of crack delamination after tensile testing. The degreeof delamination cracks significantly increased with increasing temperaturedifference (DT). Similarly, a cumulative distribution plot of pullout lengthsincreased with increasing DT. These results support the idea that the fractureresistance decreased through the extension of pre-existing microcracks onBN-rich cell boundaries, promoting the delamination cracks.

1.6 Future trends

In the last 10 years, significant advances in fibrous monolithic ceramics havebeen achieved. A variety of materials in the form of either oxide or non-oxide ceramic for cell and cell boundary have been investigated [1]. As aresult of these efforts, FMs are now commercially available from the ACR*company [28]. These FMs are fabricated by a coextrusion process. In addition,the green fiber composite can then be wound, woven, or braided into theshape of the desired component. The applications of these FMs involve solidhot gas containment tubes, rocket nozzles, body armor plates, and so forth.Such commercialization of FMs itself proves that these ceramic compositesare the most promising structural components at elevated temperatures.

Nevertheless, much work is still required to fulfill performances of FMsat room and elevated temperatures. For this goal, we address two prospectiveapproaches, from the viewpoints of micro-scale and macro-scale engineering.In micro-scale engineering, we must clearly understand the fundamentalmechanisms of crack propagation and energy absorption through FM samplesand tailor microstructures of the cell and cell boundary. Although severalfactors have been found that determine how far the deflected crack travelsdown in the cell boundary, there are still many unexamined factors affectingmechanical properties of FMs. Such close observations allow for FM designswith optimized properties (i.e., high energy absorption and high strength).Once these are accomplished, we can tailor the microstructures of the cell

*Advanced Ceramics Research, Inc., Tucson, Arizona


and cell boundary. We may reinforce the cell and cell boundary by incorporatingstrengthening or toughening phases, in the form of particulate, whisker,platelet, and fiber, into them [39–41]. In addition, a glassy phase may beincorporated into the cell boundary to enhance the densification of a weakcell boundary [42]. These approaches may be more effective in improvingmechanical properties of FMs at elevated temperatures. In addition toconventional approaches, emerging nanotechnologies may open newopportunities to design FMs in submicron-scale engineering [43]. For instance,as far as the particle size that serves as a strong phase is concerned, nano-sized particles with better control of chemical and physical characteristicscan offer potential advantages to decrease the sintering temperature and tolead to new properties. Another prospective adoption of nanotechnology isto utilize micron- or nano-sized fibers, giving an extremely large number ofcell boundaries [44].

In macro-scale engineering, we must develop innovative techniques thatare capable of producing complex FM architectures without difficulty. Todate, coextrusion and hot-pressing are commonly used to manufacture thedense FM parts. This approach may increase manufacturing cost and limitthe shape of components. To replace this conventional method, we mayadopt solid freeform fabrication (SFF) techniques to create three-dimensionalFM architectures. These SFF techniques automatically allow freeformfabrication of parts with complex geometries directly from their CAD modelsby accumulatively building three-dimensional parts layer by layer [45–47].For instance, we may employ fused deposition of ceramics (FDC) using agreen filament, prepared by conventional coextrusion [47]. To date, only afew manufacturing tools are accessible; hence, there are great opportunitiesto develop new manufacturing tools. In addition to devolvement of novelmanufacturing tools, we must exploit the new way to consolidate FMcomponents by pressureless sintering. In the case of the BN cell boundary,we may incorporate a glassy phase into the cell boundary to enhance theconsolidation of BN material. From these viewpoints, there is no doubt thatsignificant advances in FMs can be achieved in the near future and theseFMs can be used as structural ceramic components over a broad range ofapplications.

1.7 References

1. Kovar, D., King, B.H., Trice, R.W., and Halloran, J.W. (1997), ‘Fibrous monolithicceramics’, J. Am. Ceram. Soc., 80(10) 2471–2487.

2. Coblenz, W.S. (1988), ‘Fibrous monolithic ceramics and method for production’,US Patent 4 772 524.

3. Baskaran, S., Nunn, S.D., Popovic, D., and Halloran, J.W. (1993), ‘Fibrous monolithicceramics: I, Fabrication, microstructure, and indentation behavior’, J. Am. Ceram.Soc., 76(9) 2209–2216.


4. Baskaran, S., and Halloran, J.W. (1993), ‘Fibrous monolithic ceramics: II, Flexuralstrength and fracture behavior of the silicon carbide/graphite system’, J. Am. Ceram.Soc., 76(9) 2217–2224.

5. Baskaran, S., and Halloran, J.W. (1993), ‘SiC-based fibrous monolithic ceramics’,Ceramic Engineering and Science Proceedings, 14(9–10) 813–823.

6. Baskaran, S., and Halloran, J.W. (1994), ‘Fibrous monolithic ceramics: III, Mechanicalproperties and oxidation behavior of the silicon carbide/boron nitride system’, J.Am. Ceram. Soc., 77(5) 1249–1255.

7. Popovich, D., Halloran, J.W., Hilmas, G.E., Brady, G.A., Somas, S., Bard, A., andZywicki, G. (1997), ‘Process for preparing textured ceramic composites’, US Patent5 645 781.

8. Hilmas, G., Brady, A., and Halloran, J.W. (1995), ‘SiC and Si3N4 fibrous monoliths:non-brittle fracture from powder processed ceramics’, Ceram. Trans., 51, 609–614.

9. Hilmas, G., Brady, A., Abdali, U., Zywicki, G., and Halloran, J.W. (1995), ‘Fibrousmonoliths: non-brittle fracture from powder processed ceramics’, Mater. Sci. Eng.,195A, 263–268.

10. Trice, R.W., and Halloran, J.W. (1999), ‘Influence of microstructure and temperatureon the fracture energy of silicon nitride/boron nitride fibrous monolithic ceramics’,J. Am. Ceram. Soc., 82(9) 2502–2508.

11. Trice, R.W., and Halloran, J.W. (2000), ‘Elevated-temperature mechanical propertiesof silicon nitride/boron nitride fibrous monolithic ceramics’, J. Am. Ceram. Soc.,83(2) 311–316.

12. Lienard, S.Y., Kovar, D., Moon, R.J., Bowman, K.J., and Halloran, J.W. (2000),‘Texture development in Si3N4/BN fibrous monolithic ceramics’, J. Mater. Sci., 53,3365–3371.

13. Tlustochowitz, M., Singh, D., Ellingson, W.A., Goretta, K.C., Rigali N., and Sutaria,M. (2000), ‘Mechanical property characterization of multidirectional Si3N4/BN fibrousmonoliths’, Ceram. Trans., 103, 245–254.

14. Singh, D., Cruse, T.A., Hermanson, D.J., Goretta, K.C., Zok, F.W., and McNulty,J.C. (2000), ‘Mechanical response of cross-ply Si3N4/BN fibrous monoliths underuniaxial and biaxial loading’, Ceram. Eng. Sci. Proc., 21(3) 597–604.

15. Koh, Y.H., Kim, H.W., and Kim, H.E. (2004), ‘Mechanical properties of fibrousmonolithic Si3N4/BN ceramics with different cell boundary thicknesses’, J. Eur.Ceram. Soc., 24(4) 699–703.

16. Lienard, S.Y., Kovar, D., Moon, R.J., Bowman, K.J., and Halloran, J.W. (2000),‘Texture development in Si3N4/BN fibrous monolithic ceramics’, J. Mater. Sci., 35,3365–3371.

17. Harmer, M.P., Chan, H.M., and Miller, G.A. (1992), ‘Unique opportunities formicrostructural engineering with duplex and laminar ceramic composites’, J. Am.Ceram. Soc., 75(7) 1715–1728.

18. Evans, A.G. (1990), ‘Perspective on the development of high-toughness ceramics’,J. Am. Ceram. Soc., 73(2) 187–206.

19. Kerans, R.J., and Parthasarathy, T.A. (1999), ‘Crack deflection in ceramic compositesand fiber coating design criteria’, Composites, A30, 521–524.

20. Okamura, K. (1995), ‘Ceramic-Matrix Composites (CMC)’, Adv. Comp. Mat., 4(3)247–259.

21. Cook, J., and Gordon, J.E. (1964), ‘A mechanism for the control of crack propagationin all-brittle systems’, Proc. Roy. Soc. Lond., 282, 508–520.


22. Clegg, W.J., Kendall, K., Alford, N McN, Button, T.W., and Birchall, J.D. (1990), ‘Asimple way to make tough ceramics’, Nature, 357 (4 October) 455–457.

23. Trice, R.W. (1998), The Elevated Temperature Mechanical Properties of SiliconNitride/Boron Nitride Fibrous Monoliths, PhD Thesis. University of Michigan, AnnArbor, MI.

24. King, B. (1997), Influence of Architecture on the Mechanical Properties of FibrousMonolithic Ceramics, PhD Thesis. University of Michigan, Ann Arbor, MI.

25. Koh, Y.H., Kim, H.W., and Kim, H.E. (2002), ‘Mechanical properties of three-layered Si3N4/ fibrous Si3N4-BN monolith’, J. Am. Ceram. Soc., 85(11) 2840–2842.

26. Brady, G.A., Hilmas, G.E., and Halloran, J.W. (1994), ‘Forming textured ceramicsby multiple coextrusion’, Ceram. Trans., 51, 297–301.

27. Polzin, B.J., Cruse, T.A., Houston, R.L., Picciolo, J.J., Singh, D., and Goretta, K.C.(2000), ‘Fabrication and characterization of oxide fibrous monoliths produced bycoextrusion’, Ceram. Trans., 103, 237–244.

28. Advanced Ceramics Research, 3292 East Hemisphere Loop, Tucson, AZ 85705-5013, USA.

29. Koh, Y.H., Kim, H.W., Kim, H.E., and Halloran, J.W. (2004), ‘Thermal shock resistanceof fibrous monolithic Si3N4/BN ceramics’, J. Eur. Ceram. Soc., 24(8) 2339–2347.

30. Kovar, D., Thouless, M.D., and Halloran, J.W. (1998), ‘Crack deflection andpropagation in layered silicon nitride/boron nitride ceramics’, J. Am. Ceram. Soc.,81, 1004–1012.

31. McNulty, J.C., Begley, M.R., and Zok, F.W. (2001), ‘In-plane fracture resistance ofa cross-ply fibrous monolith’, J. Am. Ceram. Soc., 84, 367–375.

32. Singh, D., Goretta, K.C., Richardson, J.W. Jr., and de Arellano-López, A. (2002),‘Interfacial sliding stress in Si3N4/BN fibrous monoliths’, Scripta Mater., 46, 747–751.

33. He, M.Y., and Hutchinson, J.W. (1989), ‘Crack deflection at an interface betweendissimilar elastic materials’, Int. J. Solids Structures, 25(9) 1053–1067.

34. Koh, Y.H., Kim, H.W., Kim, H.E., and Halloran, J.W. (2002), ‘Effect of oxidation onmechanical properties of fibrous monolith Si3N4/BN at elevated temperatures inair’, J. Am. Ceram. Soc., 85(12) 3123–3125.

35. Hasselman, D.P.H. (1969), ‘Unified theory of thermal shock fracture initiation andcrack propagation in brittle ceramics’, J. Am. Ceram. Soc., 52, 600–604.

36. Hasselman, D.P.H. (1970), ‘Thermal stress resistance parameters for brittle refractoryceramics: a compendium’, Am. Ceram. Soc. Bull, 49, 1033–1037.

37. Wang, H. and Singh, R.N. (1994), ‘Thermal shock behavior of ceramics and ceramiccomposites’, Int. Mater. Rev., 39, 228–244.

38. Hirano, T. and Niihara, K. (1996), ‘Thermal shock resistance of Si3N4/SiCnanocomposites fabricated from amorphous Si–C–N precursor powders’, Mater.Lett., 26, 285–289.

39. Hai, G., Yong, H., and Wang, C.A. (1999), ‘Preparation and properties of fibrousmonolithic ceramics by in-situ synthesizing’, J. Mater. Sci., 34(10) 2455–2459.

40. Li, S.Q., Huang, Y., Wang, C.G., Luo, Y.M., Zou, L.H., and Li, C. W. (2001). ‘Creepbehavior of SiC whisker-reinforced Si3N4/BN fibrous monolithic ceramics’, J. Eur.Ceram. Soc., 21(6) 841–845.

41. Li, S.Q., Huang, Y., Luo, Y.M., Wang, C.A., and Li, C.W. (2003), ‘Thermal shockbehavior of SiC whisker reinforced Si3N4/BN fibrous monolithic ceramics’, MaterLett., 57(11) 1670–1674.

42. Trice, R.W. and Halloran, J.W. (1999), ‘Investigation of the physical and mechanical


properties of hot-pressed boron nitride/oxide ceramic composites’, J. Am. Ceram.Soc., 82(9) 2563–2565.

43. Sternitzke, M. (1997), ‘Structural ceramic nanocomposites’, J. Eur. Ceram. Soc.,17(9) 1061–1082.

44. Li, D., and Xia, Y.N. (2004), ‘Direct fabrication of composite and ceramic hollownanofibers by electrospinning’, Nano Letter, 4(5) 933–938.

45. Cawley, J.D. (1999), ‘Solid freeform fabrication of ceramics’, Curr. Op. Solid StateMater. Sci., 4, 483–489.

46. Halloran, J.W. (1999), ‘Freeform fabrication of ceramics’, Br. Ceram. Proc., 59, 17–28.

47. Rangarajan, S., Qi, G., Venkataraman, N., Safari, A., and Danforth, S.C. (2000),‘Powder processing, rheology, and mechanical properties of feedstock for fuseddeposition of Si3N4 ceramics’, J. Am. Ceram. Soc., 83(7) 1663–1669.

33

2.1 Introduction

Silicon nitride as a monolithic ceramic has earned a place in the inventory ofcommercial high-performance ceramics with excellent performance inapplications such as ball bearings, seals, cam followers, and molten metalhandling. However, like most other monolithic advanced ceramics, Si3N4

has suffered from an inherent brittleness which limits its application wheresome toughness or impact resistance is required. In response to this, workbegan in the mid-1980s to use whisker reinforcement to increase the toughnessof ceramics such as alumina and silicon nitride. The work was mainly centredaround two main whisker types – SiC and Si3N4 – until the development of‘in-situ whiskers’, or ‘self-reinforcing’ Si3N4 wherein b-Si3N4 seed crystalscould be grown into whisker-like grains during sintering. The principal workhas since been on self-reinforced Si3N4 and SiC(w)–Si3N4 composites.

An appreciable increase in toughness has been accomplished through theuse of whisker (or elongated grain) reinforcement, and K1C values of 10MPa.m1/2 are commonly reported. This chapter will summarize the materialsand fabrication processes involved in producing these materials and then goon to discuss the ensuing properties and applications. Due to various factorsincluding processing difficulties encountered with these materials and socio-economic factors, commercialization of Si3N4 composites has been somewhatlimited, but SiC(w)–Si3N4 composites are finding applications in cuttingtools and self-reinforced Si3N4 has become a popular replacement forconventional silicon nitride. Work has been continuing in this area andoptimization of processing routes to produce a viable commercial product(physically and economically) is still continuing. These routes are requiredin order to bring the enhanced properties of silicon nitride composites to therelevant industries with near-net shape capabilities and acceptable costs. Inaddition to ongoing work into the properties and behaviour of these materials,the future outlook is for more emphasis into this aspect of ‘manufacturability’and cost.

2Whisker-reinforced silicon nitride ceramics

M D P U G H, Concordia University, Canada andM B R O C H U, McGill University, Canada


2.2 Fabrication

There has been considerable progress in the processing of monolithicengineering ceramics in the last 25 years, especially in the field of sintering,microstructural control, and grain boundary refinement of silicon nitride-based ceramics. This has led to the development of gas pressure and evenpressureless sintering of silicon nitride components having excellentcharacteristics (except in the area of toughness) from green precursors preparedby a variety of techniques (die-pressing, injection moulding, slip-casting,etc.).

As with most engineering ceramics, the primary fabrication route relieson production of a green compact from Si3N4 powder (typically 40–60%dense) which is then densified by a high-temperature process such as pressure/pressureless sintering, hot-pressing, and hot isostatic pressing (HIPing). Thisdensification ipso facto results in shrinkage of the compact which producesthe major fabrication challenge for whisker-reinforced silicon nitridecomposites. This section will describe the commonly used starting materialsand fabrication routes for processing whisker-reinforced silicon nitridecomposites.

2.2.1 Materials

Silicon nitride matrix

The materials used for whisker-reinforced silicon nitride ceramics dependon two major factors: firstly the chosen processing route for fabricating thecomposite (hot-pressing, sintering, reaction-bonding, etc. – see Section 2.2.2)and secondly the properties required of the whiskers. If the silicon nitridematrix is to be formed by a sintering mechanism (hot-pressing, gas-pressureand pressureless sintering) then the starting materials will include firstlysilicon nitride powder, usually with a sub-micron size (giving a high surfacearea and hence driving force to promote sintering) and with an a-phasecrystal structure to benefit from the a Æ b transformation that occurs duringsintering, and secondly, sintering additives. These sintering additives arerequired to provide a liquid phase at high temperature through which adissolution–transport–precipitation mechanism can occur, leading todensification and the formation of a solid structure on cooling (liquid phasesintering). Unlike ceramics such as alumina and silicon carbide, silicon nitridepowders do not densify easily through solid state mechanisms. A range ofmaterials have been tried as sintering aids for silicon nitride, including magnesia,yttria, beryllia, calcium oxide and some rare earths, but common formulationsare now centred around additions of yttria and alumina totalling 10–15 wt%for sintering and lesser amounts for hot-pressing. Magnesia at levels of 4–5wt% is used for hot-pressing as lower temperatures can be used, but this will

Whisker-reinforced silicon nitride ceramics 35

also limit the maximum use temperature of the material. Unless specificallytailored to crystallize, this liquid phase forms a glassy intergranular phase oncooling which will reduce the creep resistance of the material.

Whiskers

The whiskers that are commonly added to silicon nitride to form compositesare Si3N4 and SiC (common nomenclature is to add the suffix (w) to denotewhiskers and this will be used where appropriate). Although other ceramicand metallic whiskers are available (e.g. BN, TaC, TiC, B4C and Fe), Si3N4

and SiC each have properties that make them prime candidates asreinforcements, not least of which is that they can sustain the high temperaturesand reactivity of the sintering process without being degraded.

Silicon nitride whiskers are added to improve toughness through variousmechanisms as described in Section 2.3, without introducing problems withregard to mismatch of coefficients of thermal expansion (CTE) and Young’smoduli. Since the discovery and development of self-reinforced silicon nitridein which large, elongated grains can be grown in Si3N4 during sintering toproduce an in situ, whisker-like reinforcement, ‘conventional’ silicon nitridewhiskers have not been utilized extensively as reinforcements per se but arereplaced with rod-like, b-Si3N4 seed crystals – which are, in essence, veryshort whiskers.1,2

The principal whiskers used for toughening of silicon nitride are thosebased on SiC. Silicon carbide whiskers show an excellent combination ofproperties including strength, stiffness and high temperature stability, andthey are also amenable to a variety of surface treatments, such as oxidation,which are used to augment toughening mechanisms including debondingand pullout. Silicon carbide whiskers are available in a variety of sizes,chemistries and surface conditions depending on their fabrication route. Thedimensions of these whiskers can vary but are typically of the order of 0.1–1 mm in diameter and 5–100 mm in length. The aspect ratio of whiskersranges from ª5 to ª50, but in many cases the mixing/milling processesrequired to separate whisker flocs and uniformly disperse the whiskers in thepowder mixture can result in a reduction of whisker lengths and hence aspectratio. This may not necessarily be a bad thing, however, since a reducedaspect ratio has been shown to result in improved densification, especially inpressureless and gas-pressure sintering.3

2.2.2 Fabrication

Green fabrication and sintering of compacts containing second phases posesseveral problems and when the second phase is the form of whiskers theseproblems are magnified. These problems include inhomogeneous distribution


of the second phase, thermal expansion mismatch on cooling from sintering,and most importantly whisker impingement during sintering. This last factorpresents a major obstacle in the fabrication of Si3N4 composites using sintering(gas pressure or pressureless). In order to overcome this problem there arecurrently three main approaches: application of high external pressure (hot-pressing / hot isostatic pressing (HIPing)), orientation of whiskers such thatbridging is minimal (extrusion, tape casting, etc.) and finally reaction-bondingof silicon in which no shrinkage occurs.

The first is the most commonly used in research as high uniaxial pressurecan provide ~98% or more of theoretical density.4 The major drawbacks ofthis method are the limited geometry and the preferred alignment of whiskersperpendicular to the pressing direction. This imposes limitations on theversatility and application of material processed this way. Generally, whiskers(up to 30 wt%) and powders (Si3N4 plus sintering additives) are mixed ormilled, dry or wet: if wet, the liquid is usually a solvent such as ethanolwhich is subsequently removed by drying. Hot-pressing conditions dependon the additives used but typically involve temperatures from 1650 to 1800∞Cand pressures of ª30 MPa in graphite dies surrounded by a nitrogen or argonatmosphere.5

HIPing is considerably more complex and expensive than hot-pressingbut for monolithic silicon nitride it normally produces a superior productwith near-isotropic properties at low additive contents. However, with siliconnitride whisker composites, whisker bridging and the formation of whisker‘nests’ can be more of a problem.5

In conventional (pressureless) sintering of whisker composites in an inertor nitrogen atmosphere as well as in gas-pressure sintering (ª1 MPa ofnitrogen overpressure), densification can be hampered by impingement ofwhiskers to form a skeletal network which can inhibit further shrinkage andthus result in a lower sintered density and higher percentage porosity.6 Typicallyfor more than 20 wt% whiskers the final density is less than 90% of theoreticalwhen processed this way, which leads to a reduction in many mechanicalproperties such as strength, and changes to some physical properties such asthermal conductivity.

In order to alleviate this problem several workers have attempted to orientthe whiskers in the green compact by one means or another prior to sintering.These methods include extrusion of a powder/whisker/binder mixture, slip-casting and tape-casting of slips.7–9 These techniques, although giving someimprovement, are generally limited to about 95% theoretical density (TD)after sintering, and 15–20 wt% whiskers. For higher whisker contents thereis a significant decrease in TD and flexural strength. Once again, theseprocesses result in an anisotropic structure.

The only method currently available that avoids the whole problem ofwhisker impingement is reaction bonding of silicon powder to form the


silicon nitride matrix (RBSN).10 In this process, a compact containing siliconpowder is reacted in a nitrogen atmosphere at temperatures around the meltingpoint of silicon (1300–1400∞C) for extended times (ª20 hours). Because thisnitridation generally occurs by a controlled vapour phase reaction, the increasein volume of solid material (Si Æ Si3N4) is accommodated in the pore spaceof the compact, giving a final component that has no net shape change, i.e.no shrinkage (unlike sintering or hot-pressing), but does result in a residualporosity of the order of 23%. For the manufacture of composites, whiskersare mixed with the silicon powder and the nitridation takes place in the usualway, although the kinetics may be changed, thus allowing formation of acomposite free from the problems of whisker impingement.11

Because the nitridation produces the silicon nitride matrix in situ, there isno requirement for liquid-forming sintering aids and hence no residual glassyphase unless a secondary sintering step is to be performed to reduce theresidual porosity, which then reintroduces the problem of shrinkage, albeitof a smaller magnitude.12 As the RBSN composite consists of Si3N4 directlybonded to itself and the whiskers, this material is less prone to creep atelevated temperatures but may be more susceptible to oxidation due to thehigher porosity levels. As the silicon powder compact can be formed byconventional powder processing routes (pressing, extrusion, casting), thereis a choice in whether or not the whisker reinforcement is oriented, with noeffect on the reaction bonding process.

Purely from a processing viewpoint, one can see that hot-pressing is verystraightforward and produces a good but anisotropic product, though from amanufacturing perspective this process is the least accommodating in termsof shaping. The most versatile would be the reaction bonding route exceptfor its Achilles’ heel – its high residual porosity.

2.3 Properties

One of the main strategies to improve the performance of monolithic siliconnitride ceramic is the addition of whiskers, leading to crack bridging andcrack deflection. As mentioned earlier, the whiskers most commonly used toreinforce monolithic Si3N4 are SiC and Si3N4, but more recently b-Si3N4

seeds are now used to promote in situ growth of whisker-like crystals inplace of Si3N4 whiskers. The control of the microstructure–propertiesrelationship leads to the design of materials with enhanced properties. Inaddition to the type of whisker reinforcement, the microstructure is alsomodified by the volume fraction of whiskers, their orientations and respectiveproperties and by interfacial reactions. This section describes the differentmicrostructures typically obtained for Si3N4 reinforced with Si3N4(w) orSiC(w) and the resulting changes in physical and mechanical properties.


2.3.1 Microstructure

The main mechanism reported to improve the toughness of whisker-reinforcedSi3N4 is based on crack-bridging which depends on the orientation of thefibres to the propagating crack. Components possessing whiskers orientedperpendicular to the crack propagation direction will exhibit an increase infracture strength and creep resistance at high temperature. Obviously, areduction of these properties is observed when the whisker orientation isparallel to the crack propagation direction. Therefore, the orientation of thereinforcement will dictate whether the composite has isotropic (randomorientation) or anisotropic (aligned orientation) properties.

SiC(w) reinforced Si3N4

The introduction of silicon carbide whiskers in a silicon nitride matrix iscommonly performed by mixing and densification as described in Section2.2.2. As mentioned, the sintering of silicon nitride requires liquid-formingadditives. In SiC(w)–Si3N4 composites, the reaction and wettability of thewhiskers by the sintering additives used become important parameters, asthe debonding characteristics of this interface have a strong influence onmechanical properties. Fortunately, like Si3N4, SiC possesses a native SiO2

surface layer, which is a common component of the liquid-phase sinteringprocess for silicon nitride.

Studies with different sintering additives have been performed to identifythe optimum composition. Dogan and Hawk13 have detected with TEM andXRD two scenarios regarding the crystalline state of the sintering additive:(1) crystalline Y2Si3N4O3 (millilite), b-Y2Si2O7 and an amorphous layerbridging the grains and the crystalline phase, or (2) a probably completelyamorphous yttrium–aluminium silicate. In both cases, the sintering additiveswet both the matrix and the whiskers. Unfortunately, insufficient experimentaldetails are available yet to try to demystify the wetting mechanism. High-resolution TEM of the interface for monolithic Si3N4 and Si3N4–20%SiC(w)with Y2O3 as the sintering additive has indicated a grain boundary phase ofa-Y2Si2O7 with a thin, continuous and amorphous film at the interface forthe monolithic silicon nitride. However, the addition of SiC(w) modifies thechemical composition of the sintering additive, and a glassy phase at theinterface containing a high concentration in SiO2, presumably coming fromthe surface oxide film of the whiskers, is observed.

The consolidation mechanism of many of these ceramic composites involvesa thin glassy phase linking the matrix grains, the reinforcement and thecrystalline portion of the sintering additive.14 The thermal and mechanicalproperties of all these phases are quite different, which will therefore influencethe behaviour of the composite. Although the coefficient of thermal expansion


(CTE) of the crystalline phase would change considerably depending on theinitial sintering additive added, the Y2Si3N4O3 is a member of the gehlenitefamily, which possesses a CTE of around 1–2 ¥ 10–6/∞C15 as opposed to5 ¥ 10–6/oC for the amorphous phase.16 Knowing that the CTE of Si3N4 isaround 3.2 ¥ 10–6/oC and that of SiC is 4.5 ¥ 10–6/oC, the volume fractionand composition of the different phases present will lead to complex interfacialresidual stress patterns. For example, the presence of a crystalline phase withsuch a low CTE will lead to compressive residual stresses at room temperaturefor the crystalline boundary phase. The opposite will be observed for theamorphous boundary phase where tensile residual stresses will be present.This change in residual stress will influence the mechanical properties, asone scenario will promote better toughening and the other better strength.

The morphology of the added whiskers also plays an important role duringprocessing as does the actual mixing technique. Whiskers with high aspectratio can have a deleterious effect on green density consolidation andsinterability, decreasing the fracture toughness of the composite.3 Anotherimportant modification of the processing parameters by the presence of SiC(w)is that of the kinetics of the a Æ b transformation. This transformation iswell known in monolithic Si3N4 ceramics to improve the fracture toughnessby the development of elongated b-grains. The presence of SiC(w) retardsthe a Æ b phase transformation, thus reducing this beneficial effect.

Si3N4(w) reinforced Si3N4

As mentioned earlier, this type of microstructure can be obtained by addingsilicon nitride whiskers to the Si3N4 powder as for SiC whisker compositesor more commonly by the growth of large and elongated b-grains from rod-like b-seed crystals in an otherwise equiaxed b-Si3N4 matrix during sintering.These high-aspect-ratio grains act similarly to whiskers. Therefore, controlof the processing parameters (mixture and composition of the raw materials,hot-pressing vs. liquid phase sintering, sintering additives, etc.) directly affectsthe composite microstructure and therefore its response to service stresses.The growth of the elongated b-grains is anisotropic; a higher growth rateoccurs along the c-axis of the grain and is also influenced by the presenceand composition of the sintering additive.

The control of the orientation of the resulting elongated grain is performedby orienting the initial b-grain seeds, which grow at the expense of the smallmatrix grains. Therefore, their initial alignment leads to an orientedmicrostructure, and the most favourable processes to produce an orientedseed structure are tape casting and extrusion. The growth of the Si3N4 whiskersin an oriented green body is such that they have a lower probability ofimpeding each other, resulting in a microstructure with a reinforcement phasepossessing a higher aspect ratio than for the randomly oriented seed mixtures.


Lee et al.17 have reported that the growth of elongated b-Si3N4 whiskers isa function of the initial starting powder size and occurs by the dissolution ofsurrounding grains. Whiskers with a higher aspect ratio are obtained whenfiner seeds are used, and a critical particle size is important as 2-D nucleationand growth must prevail, since larger particles will promote the growth of amore equiaxed microstructure. Obviously, the sintering parameters modifythis behaviour; whisker-containing microstructures will be obtained at highertemperatures when larger seed particle sizes are used.

Control of the microstructure (in this case, primarily the volume fractionof whiskers) is usually achieved through the initial ratio of a/b. As mentionedearlier, the addition of b-seeds in an a-powder produces the elongated b-grains in a fine b-matrix (subsequent to the a Æ b transformation) accompaniedby a residual glassy phase. In addition to growth, these b-seeds act as nucleatingagents for the a Æ b transformation. However, the existence of a criticalamount of nuclei, 10 wt%, was proposed by Emoto and Mitomo18 abovewhich a reduction of the growth driving force occurs. Therefore, the generationof a distinct bimodal microstructure, especially one with large, elongatedgrains, well dispersed in a fine, submicron-grain-size matrix, is an importantfactor in optimizing the fracture resistance and the fracture strength of thematerial.

However, the presence of the whisker reinforcement is not the only factoraffecting the response to mechanical stresses. Silicon nitride is commonlyproduced through liquid phase sintering, and therefore a residual quantity ofsintering additive phases (crystalline or amorphous) will remain at the grainboundaries, and their respective volume fractions will modify the tougheningmechanisms. In addition, the chemical composition of the sintering additivewill influence the grain growth rate during processing, and different grainmorphologies and sizes can be obtained.

The use of a sintering additive containing a high concentration in aluminaand nitrogen has been shown to favour the development of a distinct epitaxialb¢-SiAlON layer at the interface between the Si3N4 whisker and the oxynitrideglass.19 Interfacial debonding was found to vary with the glass composition,where for low yttrium–aluminium ratios, the b¢-SiAlON layer grew on theSi3N4 whiskers and the b¢-SiAlON/oxynitride glass interface was reported tobe more resistant to debonding as the chemical gradients are more gradual.Therefore, the mechanical properties of the composite are strongly influencedby the z value (concentration of substituted Al and O) of the SiAlON near theb¢-SiAlON/glass interface. Experiments have shown that for sintering additiveswith lower z value, an improvement of toughness of 30% was obtained.19

The absence of the b¢-SiAlON layer was found to improve the fractureresistance of the composite by debonding, crack deflection and whiskerpullout.20 However, the concentration in Al in the b¢-SiAlON layer is controllednot only by the sintering additive but also by the extent of grain growth. The


processing conditions used to establish the interface, such as temperatureand time, will modify the interfacial chemistry by the control of grain growth.Therefore, the use of Al2O3 as sintering additive should not be proscribed, asalumina plays an important role in the densification and a low quantity willenable a limited growth of b¢-SiAlON layer, which will still have low interfacialstrength and promote the targeted debonding behaviour that leads toimprovement in toughness.

As opposed to composites reinforced with SiC(w), the residual stressesformed at the interface should be lower as the CTE of the matrix and thereinforcement phase are similar. Several modeling studies have shown thatthe composition of the grain/glass interface has a negligible effect on themagnitude of the residual stresses and therefore has no significant influenceon the fracture behaviour.19

2.3.2 Improved properties

In whisker-reinforced ceramic composites, the modification of the mechanicalproperties is dictated by the strength at the grain/grain boundary phase interfacesas discussed above. This alteration of the mechanical properties compared tothe monolithic Si3N4 is observed on a macro-scale through several mechanisms:whisker debonding, whisker pullout, crack bridging and/or crack deflection.However, the common idea in these three toughening mechanisms is toincrease the energy needed to extend a crack. A brief description of thesemechanisms will be presented to help understand their impact on mechanicalproperties.

The concept of crack deflection arises from the lower-strength grainboundary present in polycrystalline materials and is governed by the whiskershape and volume fraction, although elastic moduli and thermal expansioncoefficient have a role. A propagating intergranular crack will have a tendencyto follow the easiest path and therefore the crack orientation may change asthe crack propagates. This change in orientation reduces the average stressintensity at the crack tip, because the crack propagation plane is no longernormal to the applied tensile stress. This toughening mechanism is wellaccepted by the community for explaining the lower toughness of singlecrystal versus polycrystalline ceramics.

Crack bridging implies the mechanical support by a second phase to helpreduce the opening of a propagating crack. The second phase acts as aligament behind the crack front, reducing the stress intensity factor. Thestress supported by the ligament increases slowly with distance behind thecrack tip, and greater crack-opening displacements are achieved in the bridgingzone, enhancing the fracture toughness. Crack bridging is often enhanced byincreasing the volume of the reinforcement phase, the ratio of the elasticmodulus of the ceramic over that of the fibre, Ec/Ef, and the ratio of the


fracture energy of the bridging ligament to that of the reinforcement/matrixinterface.

Crack bridging is also complemented by the contribution of pullout of thefailed whisker reinforcements. The pullout operation consumes energy whichwould otherwise contribute to the advancement of the crack front and thusenhances toughness.

However, these strengthening mechanisms are not the answer to all theproblems. The improvement of the mechanical properties by the addition ofwhiskers greatly influences the long-crack toughness but has little or noinfluence on the short-crack toughness. In the design of a component fromwhisker-reinforced ceramics, the geometry and size of the final part influencesthe desired microstructure. Because of the difficulties in processinghomogeneous, large-scale ceramic components, long-crack toughness isdesirable in this case as it increases the tolerance to flaws (e.g. microstructuralheterogeneities, such as weak grain boundaries, large grains and residualstresses) and improves the service life. The toughness of this compositeincreases with the crack length due to internal friction resistance. On theother hand, the design of small-sized ceramic composites requires a morehomogeneous microstructure – the fracture toughness is dictated by theheterogeneity in the microstructure, which should be reduced as much aspossible in small pieces. In this case, the toughness becomes more independentof the microstructure as crack bridging due to internal friction resistance isless efficient. Therefore, more optimization of the microstructure must bedone to improve the resistance to short-crack propagation. The followingsections summarize the improvements of physical and mechanical propertiesobtained for whisker-reinforced silicon nitride composites to date.

Physical properties

Coefficient of thermal expansion

The coefficient of thermal expansion (CTE) of composite materials usuallyfollows the simple rule of mixtures (or more complex models), based on theCTE of the respective components, their volume fraction and the volumefraction of interfacial phases. Based on these models, a Si3N4–Si3N4(w)composite should possess a similar CTE to monolithic Si3N4 ceramic (3.2 ¥10–6/oC). Obviously, the chemical composition of the sintering additive willhave a certain influence but should remain within the variations observed formonolithic Si3N4.

However, in the case of SiC(w), the CTE of the reinforcement phase (SiC:4.5 ¥ 10–6/oC) is higher than that of the composite matrix, implying an increasein the composite CTE when whiskers are added to silicon nitride. Jia et al.21

reported that the CTE of 30 vol% SiC(w)–Si3N4 reaches nearly 7 ¥ 10–6/oC,


which is more than the double the value of monolithic Si3N4. However, theirmeasurements were performed on samples with densities varying between86–91% and with residual a-phase levels reaching up to 10%, both of whichcan affect the CTE values.

Thermal conductivity

The thermal conductivity of a Si3N4 composite is controlled by the level ofporosity, the b-phase content, the amount and conductivity of the reinforcementphase (Si3N4 vs. SiC), the amount of glassy phase and any orientation effect.In Si3N4–Si3N4(w), the thermal conductivity will increase with the densityand with the amount of b-phase, as it possesses a higher thermal conductivitythan a-phase. In addition, orientation of the whiskers produces an anisotropicthermal conductivity behaviour, where a higher thermal conductivity, up to1.5 times higher in fact, is observed in the direction of orientation comparedwith the direction perpendicular to the grain orientation.22 Therefore, long b-grains with fewer triple junctions are favourable for conductivity. Finally,the volume fraction of the glassy phase influences conductivity, as the sinteringadditives typically used possess a lower thermal conductivity than the matrixand act as a thermal barrier.

The prediction of the thermal conductivity is not as simple for the compositereinforced with SiC whiskers. In fact, as mentioned earlier, the thermalconductivity of the reinforcement phase plays an important role. However, forSiC whiskers, their chemical composition can vary drastically depending onthe manufacturing process, which is not the case for Si3N4 whiskers. In fact,the influence of the manufacturing process is drastic: the thermal conductivityof SiC(w) produced by the vapour–solid process is around 20 W/m K asopposed to 100–250 W/m K for whiskers produced by the vapour–liquid–solid process. Using a Si3N4 matrix with a nominal thermal conductivity of48.5 W/m K, the thermal conductivity of the Si3N4–SiC(w) composite canincrease or decrease depending on the type of SiC whiskers added.23

Mechanical properties

Flexural strength

The targeted properties in Si3N4 composites are high strength and fracturetoughness for high-temperature applications. Figure 2.1 presents four-pointflexural strength values for hot-pressed Si3N4–Si3N4(w) composites withdifferent whisker volume fractions as a function of temperature (MgO assintering aid).24 The strength for a monolithic Si3N4 is added for comparisonpurposes. The results demonstrate that the addition of whiskers improves theflexural strength at low and intermediate temperature (approximately 1000oC)


but becomes less significant at higher temperature (1400oC). The addition ofsome whiskers can improve strength properties but additions over a criticalamount will have a deleterious effect on the microstructure. The presence ofmore than 15 vol% of Si3N4 whiskers reduces the elongated grain growth ofthe Si3N4 matrix, thus reducing the desired effect of adding whisker.25 Thiscan be observed by the similarity in the curves for the composites possessing0, 20 and 35 vol% whiskers. By increasing the test temperature, thestrengthening mechanism of the composite changes, passing from crackbridging and whisker pullout to a softening of the grain boundary and stressrelaxation at high temperature, resulting in easier whisker pullout. The changein failure mechanism is also apparent by the change in the fracture surface,where round, smoothed surfaces and grain boundary cavities are observedfor samples tested at 1400oC. As mentioned in Section 2.2 the use of MgOas sintering additive is known to reduce the high-temperature strength of thecomposites compared to Y2O3–Al2O3 additives. Therefore, higher flexuralstrength at higher temperature can be expected for composites sintered withyttria–alumina mixtures.

The influences of the Si3N4 matrix grain size distribution and Si3N4(w)aspect ratio on the flexural strength at room temperature have been reportedby Becher et al.26 and are presented in Table 2.1. The results show thatimproved strength is obtained for a distinct bimodal grain size distributionand that the characteristics of the raw materials are very important. In suchcases, the mechanical behaviour of the composite is getting closer to theoptimum microstructure for crack bridging and whisker pullout. Anotherfactor that appears to affect strength is the ratio of yttria to alumina in thesintering additive which changes the phases present at the interface; thepresence of a b¢-SiAlON layer reduces the flexural strength.19

0 300 600 900 1200 1500Temperature (∞C)

0% Si3N4w 5% Si3N4w20% Si3N4w 35% Si3N4w

1000

800

600

400

200

0

Flex

ura

l st

ren

gth

(M

Pa)

2.1 Flexural strengths as a function of temperature for silicon nitridecomposites.24


As opposed to the processing of Si3N4–Si3N4(w) composites, where themajority of whisker formation occurs during sintering by growth of b-Si3N4

seeds, the processing of composites reinforced with SiC(w) is harder due tothe impingement of the whiskers, which can lead to lower final densities thatin turn also affect the composite properties. Figure 2.2 presents a summaryof the room-temperature flexural strength of some Si3N4–SiC(w) composites.For hot-pressed composites, it appears that the presence of SiC(w) reducesthe strength and there is no sign of an optimum volume fraction ofreinforcement, as in Si3N4–Si3N4(w). In these two cases, the sintering additiveswere similar (a mixture of yttria and alumina) but their ratio differs. Theflexural strength results correlate with the final density and with the chemicalphases present at the grain boundaries. As noted earlier, improved flexuralstrength is obtained with sintering additives comprising a high ratio of Y2O3

to Al2O3 combined with high density. For composites produced by a combinedreaction-bonding/sintering route where there is less shrinkage, higher flexuralstrengths are obtained for 5% compared to 15% SiC(w).27

High-temperature four-point bending tests of Si3N4 and Si3N4/20%SiC(w)composites prepared with no external sintering aids by hot isostatic pressing

Table 2.1 Fracture strength as a function of grain size and distribution26

Average grain size Type of distribution Fracture strength (MPa)

0.4 mm Monomodal 660 ± 1651 mm Broad 850 ± 750.2 mm/0.5 mm Bimodal 925 ± 750.3 mm/2 mm Bimodal 1144 ± 126

0 5 10 15 20 25 30Volume fraction SiC(w) (%)

1000

900

800

700

600

500

400

300

200

100

0

Flex

ura

l st

ren

gth

(M

Pa)

Process Ratio Y2O3/Al2O3 Ref

Reaction bonded – hot-pressed 8 [27]Hot-pressed 1.5 [21]Hot-pressed 2.66 [28]

2.2 Flexural strength as a function of SiC whisker content.


show no reduction in strength between room temperature and 1400oC forboth the monolithic ceramic and the composite, demonstrating the importanceof the additives and grain boundary softening on the high-temperature failuremechanism.29 However, the strength values were limited to ª500 MPa, whichis half the room temperature strength obtained when sintering additives areused in the processing.

Young’s modulus

It is well known that the Young’s modulus of a composite can be calculatedby the rule of mixtures for long-fibre reinforced material. In the case ofwhiskers, the rule of mixture is also applied to estimate the change of modulus(conventionally, reinforcements are added to improve the stiffness of a material,though for ceramic matrix composites this is not always the primary concern).

For Si3N4–Si3N4(w) composites one would not expect any significantchange in elastic modulus; however, as 100% solid density is rarely achieved,there may be some decrease depending on the level of porosity present.

For SiC(w) reinforced composites, due to the higher Young’s modulus ofthe SiC whiskers (ª420 GPa) compared to monolithic silicon nitride (320GPa), an increase in stiffness should be observed, and for Si3N4–20vol%SiC(w)composites produced by extrusion, reaction bonding and hot isostatic pressing,the highest value observed was 350 GPa which is actually higher than thatpredicted by the rule of mixtures.14 The reason for this is most likely in theestimation of the elastic modulus of the SiC whiskers.

Toughness

There are two main techniques used to measure the fracture toughness ofceramics: fracture stress and hardness indentation. The former measures theload to fracture of a pre-cracked specimen using a single edge notched beam(SENB) or a chevron notched beam (CNB) sample. The main drawback ofthis technique is ensuring that the crack tip is atomically sharp. The secondmethod uses the crack formed at the corners of the indentation producedduring a Vickers indentation hardness test. This technique is rapid and relativelyinexpensive. However, the toughness values measured are those of the surface,unlike the values obtained by fracture of the pre-cracked beams which are ameasure of the bulk material properties.

The main tool for comparing the toughness of monolithic and whisker-reinforced ceramics is through R-curve behaviour. As mentioned previously,the presence of whiskers can result in reduced crack opening and thereforethe toughness becomes dependent on the propagating crack length, as opposedto the typical Griffith theory in which the initial crack or flaw is assumed topropagate instantaneously and completely. Figure 2.3 presents a schematic


representation of a propagating crack in (a) a monolithic ceramic, and (b) aceramic composite reinforced with whiskers, and also a representative viewof the influence of R-curve behaviour on (c) fracture toughness and (d)strength. In essence, ceramics that exhibit R-curve behaviour are more tolerantto the presence of flaws than those which do not.

Si3N4-Si3N4(w) compositesAs opposed to composites reinforced with SiC(w), the formation of theelongated microstructure in the Si3N4–Si3N4(w) system is occurring in situ,during the consolidation operation. Several researchers have reported thatthe main factor affecting the toughness of these composite is the volumefraction of reinforcement, but the presence of a bimodal grain distributionwas also found to have a major role in the toughening mechanism. Figure 2.4nicely summarizes the microstructure evolution vs. toughening relationoccurring for seeded and non-seeded silicon nitride composites.30 MonolithicSi3N4 with a fine equiaxed matrix shows a modest R-curve response with asteady state toughness of only 3.5 MPa.m1/2. Longer hot-pressing times orhigher pressing temperatures will enhance the formation of elongated grainsand will be accompanied by an increase in R-curve behaviour. However,

Flaw size(c)

Flaw size(d)

Str

eng

th

Tou

gh

nes

s R-curve behaviour

No R-curve behaviour

No R-curve behaviour

R-curve behaviour

(a) (b)

c c

a1 a2

2.3 Schematic representation of cracking in (a) a monolithic ceramicand (b) a ceramic composite reinforced with whiskers (a1 > a2);(c) influence of R-curve behaviour on fracture toughess, and (d)influence on strength, in ceramics and composites.


materials with a broad grain diameter distribution and with a large fractionof elongated grains exhibit only a modest increase in toughness combinedwith a lower strength. The control of the microstructure by the addition of b-seeds shows that a major improvement in the R-curve behaviour can beachieved and fracture toughness values of 10 MPa.m1/2 have been reported.26

The microstructure developed in such cases comprises large elongated grainsembedded in a fine matrix.

From a microstructural aspect, increasing the size of the reinforcing grainswill enhance contributions from both frictional bridging and pullout. In addition,enhanced interfacial debonding, number of bridging reinforcements, andlength of the debonded interfaces should all result in a more rapid rise in theR-curve, and therefore the toughness. However, the control of the initialcomposition is very important, as an overabundance of seeds will lead tograin impingement and will reduce their growth and consequently inhibit theformation of the bimodal microstructure. A surplus of seeds in conjunctionwith poor processing will favour the formation of clusters of large grains,which are known to reduce the strength and toughness.

As mentioned earlier, the toughness of the material is a function of theorientation of the whiskers, where isotropic properties will be obtained forrandomly oriented whiskers and anisotropic properties for orientedmicrostructures. The influence of whisker orientation as well as the ratio ofyttria to alumina on the toughness of Si3N4–Si3N4(w) composites as measuredby Sun et al.20 are presented in Table 2.2. In all cases, the fracture toughnessof the composite is higher in the direction normal to the whiskers, whichconfirms that more energy is required for a crack to propagate throughwhiskers by crack deflection and crack bridging than it is along the whisker/

0 300 600 900 1200 1500Crack length (mm)

12

10

8

6

4

2

0

Frac

ture

res

ista

nce

(M

Pa.

m1/

2 )

Gas pressured sintered – seeded

Hot-pressed 2 h; unseeded

Hot-pressed

Hot-pressed 0.33 h; unseeded

2.4 Effect of processing conditions on R-curve behaviour for in-situSi3N4 composites.26


matrix interface. As discussed earlier in the microstructure section, theconcentration and chemistry of the sintering agent can change the interfacialreaction, and for low Y2O3 to Al2O3 ratios, a b¢-SiAlON/oxynitride glasslayer is formed, giving a much stronger matrix/whisker interface, leading toa reduction in the efficiency of the toughening mechanisms.

Table 2.3 presents the variation in toughness of gas-pressured, sintered,seeded Si3N4–Si3N4(w) as a function of sintering time. The results clearlydemonstrate that the toughness evolves with the microstructure. The kineticsof formation of the elongated grain population is a function of severalparameters such as time, temperature and volume of seeds. The high-toughnesssamples can also be obtained faster by using higher sintering temperatures,but other problems may occur such as low cost-effectiveness, high residualstresses or interfacial reactions. The toughness values reach a peak after 90–180 minutes of sintering time, which is in accordance with Faber’s theorystating that the crack deflection mechanism becomes independent of volumefraction above 20%.31 For smaller volume fractions, the toughness of thecomposite is proportional to the mean grain diameter. The reduction oftoughness for the composites sintered for 9 hours is associated with interfacialchemistry – possibly increased crystallization of the grain boundary phase orincreased thickness of the interfacial reaction zone.

Si3N4–SiC(w) compositesOne of the main objectives in adding SiC(w) to Si3N4 matrices is to improvethe fracture toughness. Figure 2.5 shows a crack propagating through areaction-bonded silicon nitride reinforced with SiC whiskers which has beendeflected by a whisker, resulting in an increase in absorbed energy. This

Table 2.2 Dependence of fracture toughness on grain orientation as afunction of sintering additive20

Sintering additive Toughness perpendicular Toughness parallelto whisker direction to whisker direction(MPa) (MPa)

4.0Y2O3–2.8Al2O3 7.5 ± 0.2 6.9 ± 0.25.0Y2O3–2.0Al2O3 8.7 ± 0.4 7.8 ± 0.36.25Y2O3–1.0Al2O3 10.6 ± 0.2 9.2 ± 0.2

Table 2.3 Variation in toughness as a function of sintering time32

Sinterig 10 45 90 180 540time (min)

Toughness 6.7 ± 0.2 7.6 ± 0.2 7.9 ± 0.3 8.1 ± 0.3 6.9 ± 0.2(MPa.m1/2)


toughening behaviour rests on numerous variables including characteristicsof the SiC whiskers and their surfaces. There are several manufacturingprocesses available to produce SiC whiskers, which result in whiskers, andhence composites, possessing a wide range of properties. For example,Rossignol et al.5 have demonstrated that Si3N4–30%SiC(w) compositesmanufactured with the same procedure but with differing SiC whiskers mayhave nearly 20% difference in indentation fracture toughness (6.8 ± 0.9 vs.8.2 ± 0.6 MPa.m1/2) for samples with similar porosity levels (98.7 vs. 99.7%TD).

As mentioned earlier, fracture toughness values can vary significantly dueto the method of measurement, the reinforcement properties, manufacturingdefects, etc. Taking into account these possible factors, a small sample of thefracture toughness values reported in the literature for Si3N4–SiC(w) compositeswith different volume fractions and different sintering additives is presentedin Fig. 2.6. The fracture toughness values of the samples reaction-bondedand then sintered with yttria and alumina are high even in the monolithicstate and the sample do not show any major improvement of toughness withaddition of SiC(w). In contrast, the composites sintered with the AlN-rareearth additives possess a lower toughness for the monolithic sample whichimproves with the addition of whiskers. As mentioned earlier, two majorfactors influence the toughening mechanisms: (1) the interfacial reactionlayer, which affects the debonding and pullout, and (2) the bimodal grainsize distribution for crack deflection. For the sample reaction-bonded andsintered with Y2O3–Al2O3 the microstructure revealed, in addition to the

2.5 Crack deflection in a SiC(w)-reinforced reaction bonded siliconnitride.


SiC(w), the presence of Si3N4 needles in an equiaxed Si3N4 matrix. Thetoughening is attributed to these phases and thus is not significantly affectedby the presence of SiC(w). Regarding the samples sintered with the AlN-based sintering additive, the presence of elongated Si3N4 grains was notreported and the improvement of toughness is directly related to the presenceof SiC(w). These results also reinforce the necessity of having higher Y2O3

concentrations in order to improve toughness, as better interfacial reactionsoccur at the interface with the Si3N4 grains. Finally, the results demonstratethat composites reinforced with 20 vol% of SiC(w) can reach similar valuesof toughness to microstructurally optimized Si3N4–Si3N4(w) composites.

Figure 2.7 summarizes the variation of toughness (indentation method) ofSi3N4 composites as a function of volume fraction of SiC whiskers possessingdifferent aspect ratios (R). In all cases, similar sintering additives were used.The maximum fracture toughness, approximately 10 MPa.m1/2 for 20 vol%SiC(w), is reached for the largest aspect ratio (15). The moderate improvementfor whiskers with lower aspect ratio is attributed to their smaller contributionto the main toughening mechanisms: crack deflection and crack bridging.The rapid decrease in toughness of the composites with the highest aspectratio is associated with the reduced density of the composite. In fact, suchlong whiskers impede consolidation and densities of only 2.5 g/cm3 wereobtained. Therefore, the optimum whisker aspect ratio for producing acomposite with high fracture toughness should be kept at around 15, whichalso allows a respectable level of density.

Notch-beam technique

Vickers indentation technique

0 2 4 6 8 10 12 14 16 18 20Volume fraction SiC(w) (%)

11

10

9

8

7

6

5

4

3

2

1

0

Frac

ture

to

ug

hn

ess

(MP

a.m

1/2 )

Process Sintering additives Ref

Rx bonded – hot pressed 8wt% Y2O3–1wt% Al2O3 [27]Hot-pressed 5wt% AlN–5wt% CRE* [33]Hot-pressed 5wt% AlN–2.5wt% CRE [33]

2.6 Fracture toughness as a function of SiC whisker content fordifferent composites.


Microhardness

The microhardness values reported for silicon nitride reinforced with SiCwhiskers are somewhat contradictory when compared to unreinforced Si3N4;sometimes they are higher and sometimes lower. A main factor affecting thehardness is the presence, chemical content and composition of the sinteringadditive, which dictates the bonding characteristics between the reinforcementand the matrix. Dogan and Hawk13 reported a higher increase of microhardnessover the monolithic material for their Si3N4–20%SiC(w) composite withcrystalline sintering additives (19.0 vs. 15.6 GPa) than for their compositewith an amorphous phase boundary (16.5 vs. 15.0 GPa). Baldacim et al.35

have measured the variation of the microhardness as a function of the sinteringadditive (mixtures of AlN and Y2O3) and volume fraction of SiC(w) (10, 20and 30 vol%). Their results show that the microhardness decreases with thevolume fraction of reinforcement and, in addition, the measured values arehigher for the sintering additive with a lower AlN/Y2O3 ratio. They correlatethis higher hardness with the presence of crystalline Y2Si3N4O3, not detectedfor the higher ratio of AlN/Y2O3.

Wear resistance

In monolithic or ceramic matrix composites, increases in hardness or toughnessare not the only factors that lead to improved wear resistance. Microstructure

2.7 Effect of aspect ratio R on fracture toughness of SiC(w)–Si3N4composites.

0 5 10 15 20 25 30Volume fraction SiC(w) (%)

R = 15

R = 25R = 5

Process Sintering additives Ref.

Hot-pressing Y2O3–Al2O3 [34]Pressureless sintered 8wt% Y2O3–2wt% Al2O3 [3]Pressureless sintered 8wt% Y2O3–2wt% Al2O3 [3]

10

9

8

7

6

5

4

3

2

1

0

Frac

ture

to

ug

hn

ess

(MP

a.m

1/2 )


is also a major factor in the equation. The addition of a second phase, whiskersfor instance, creates internal stresses. The difference in CTE in multiphasecomponents, as well as anisotropic thermal expansion in single-phase materials,form tensile and compressive residual stresses at phase boundaries.

As seen previously, the addition of silicon carbide can increase the toughnessand the hardness of silicon nitride composites, though the wear behaviourmay be decreased. As described above, some sintering additives result inlower wettability of the whiskers, thus giving a lower interfacial strength,leading to debonding of the whiskers which, although increasing toughness,increases the wear rate by surface pullout compared to monolithic Si3N4.13

Thermal shock resistance

Thermal shock is one of the main drawbacks in the utilization of ceramicsfor high-temperature applications, and one of the aims of making CMCs isto improve thermal shock resistance. The thermal shock of ceramic materialsis influenced by many factors such as strength, Young’s modulus, fracturetoughness, thermal conductivity and thermal expansion coefficient. As whisker-reinforced silicon nitrides can show an improvement in room-temperaturestrength and fracture toughness, one would expect that these CMCs will alsopossess increased thermal shock resistance. Jia et al.21 have reported that theaddition of SiC whiskers improves the thermal shock resistance to catastrophicfailure but decreases the resistance to fracture initiation, demonstrating thatthe situation is not straightforward.

Creep resistance

Creep resistance for both types of whisker reinforcement in sintered compositesis governed by grain boundary sliding via a viscous flow mechanism occurringalong the amorphous phase. Depending on the temperature, the volume fractionand the viscosity of the glass system, if the rigidity of the residual sinteringadditive phase can be maintained no significant creep is observed. Nixon etal.36 have reported that creep of hot-pressed Si3N4–20%SiC(w) with Y2O3

and Al2O3 as additives is negligible below 1300∞C in a nitrogen environmentas the grain boundaries are sufficiently rigid and that relatively littledisplacement along the grain boundaries is observed. Similar results havereported similar creep results for Si3N4 reinforced with 0–20 vol% SiC(w)with MgO additive.37 In both cases, the researchers reported that the volumefraction of the reinforcement had no effect, showing that the viscosity of theglassy phase is the main active mechanism for creep. Above 1300oC, cavitationat Si3N4–Si3N4 grain boundaries becomes more important as the viscousflow of the sintering additives is increased;38 however, no such cavities wereobserved at the SiC–Si3N4 boundaries. In some cases, 20 vol% SiC(w)


composites show higher creep rates than the unreinforced Si3N4 and this isattributed to a higher resistance to devitrification of the grain boundaryphase due to the SiO2 present at the surface of the SiC whiskers. For a 30vol% SiC(w)–Si3N4 composite (MgO additive), the creep resistance in airwas improved by the whiskers, which may suggest that the higher volumefraction of reinforcement may play a more important role in improving creepresistance as more mechanical impingement should occur, reducingdeformation.39 This is contrary to the behaviour of monolithic ceramicssintered with MgO additives: they exhibit lower creep resistance as theamorphous Mg2SiO4 glass composition is known to have a lower viscositythan Y2O3–Al2O3 base glass for the same test temperature, since Y2O3 isconsidered as a devitrifying agent in this glass composition.

Fatigue resistance

The crack growth rate during cyclic fatigue of materials is commonly recognizedto be lower than for a sustained load test. Whisker-reinforced silicon nitridecomposites are no exception to this rule. However, the presence of the secondphase (whiskers) and sintering additive, depending on the processing route,will also influence the fatigue resistance. Zhu et al.40 have reported that forgas-pressure sintered Si3N4–20vol%SiC(w), the static and cyclic fatigue lifeare equal for temperatures up to 1000oC but a decrease of the static strengthwas observed at higher temperature. This change in properties is related tothe change of fracture mode which is related to the microstructure. Attemperatures up to 1000oC, the fracture propagation occurs in a mixture ofintergranular and transgranular modes, whereas nucleation, growth andinterlinkage of cavities at the front of the propagating crack is observed athigher temperatures.

2.4 Applications

As discussed previously, ceramic matrix composites were originally developedto overcome the brittleness of monolithic ceramics. Thermal shock, impactand creep resistance can also be improved, making CMCs premium replacementchoices for some technical ceramics. Industrial applications such as inautomotive gas turbines or advanced cutting tools are already taking advantageof such characteristics.

2.4.1 High-temperature gas turbines

Between 1990 and 1997, the Japanese successfully developed a 100 kWautomotive ceramic gas turbine in which the severe operating conditionsrequired high-performance materials. 41 Among the different CMCs developed,


the turbine rotor (rotational component) and the back plate (stationarycomponent) were made of SiC(w)–SiAlON or Si3N4(w)–Si3N4 ceramics.

The turbine rotor, made of in situ reinforced Si3N4 ceramic, exhibits aWeibull modulus of over 20, indicating the reliability of the component. Thispart has confirmed its potential through hot spin tests at 1200oC (70,000rpm), a flexural strength at 1200oC of 960 MPa, a fracture toughness of7 MPa.m1/2 and an oxidation resistance at 1200oC for 200 hours (890 MPa).Regarding the back plate, both SiC(w)–SiAlON and Si3N4(w)–Si3N4

compositions sustained the stationary test (26 and 31 hours respectively at1350oC, 5 atm), proving their suitability for such components.

2.4.2 Cutting tools

The trend in machining is to raise throughput, i.e. increase machining speedand material removal rate. The application of CMCs becomes eminentlyfavourable under such conditions. Silicon nitride ceramic (sintered withadditive) is conventionally used to mill or rough-turn cast irons. Self-reinforcedSi3N4(w)–Si3N4 ceramics, produced by sintering, are showing improved high-temperature properties and excellent fracture toughness compared to theconventional Si3N4 as no softening of the glass phase occurs at hightemperature.42 A number of suppliers are using this fabrication method forcutting tool inserts.

SiC(w)–Si3N4 ceramics are also used for turning hard parts (hard steels oriron-based parts with hardnesses of 48–64 HRC). Turning has been shown togive higher material removal rates compared to conventional grindingoperations for these materials, but this process generates higher temperaturesand cutting forces, justifying the utilization of CMCs. Unfortunately, thesilicon nitride matrix possesses a lower resistance to chemical wear arisingfrom reaction between the materials being machined and the tool tip, limitingtheir uses.43

2.5 References

1. Tani, E., Umebayashi, S., Kishi, K. and Kobayashi, K. ‘Gas-pressure sintering ofSi3N4 with concurrent addition of Al2O3 and 5 wt% rare earth oxide: high fracturetoughness Si3N4 with fiber-like structure’, Am. Ceram. Soc. Bull., 65[9] (1986)1311–1315.

2. Li, C.W. and Yamanis, J. ‘Super-tough silicon nitride with R-curve behaviour’,Ceram. Eng. Sci. Proc., 10[7–8] (1989) 632–645.

3. Sneary, P.R., Yeh, Z. and Crimp, M.J. ‘Effect of whisker aspect ratio on the densityand fracture toughness of SiC whisker reinforced Si3N4’, J. Mat. Sci., 36 (2001)2529–2534.

4. Shalek, P.D., Petrovic, J.J., Hurley, G.F. and Gac, F.D. ‘Hot-pressed SiC whisker/Si3N4 matrix composites’, Am. Ceram. Soc. Bull., 65[2] (1986) 351–356.


5. Rossignol, F., Goursat, P. and Besson, J.L. ‘Microstructure and mechanical behaviourof self-reinforced Si3N4 and Si3N4–SiC whisker composites’, J. Eur. Ceram. Soc., 13(1994) 299–312.

6. Tiegs, T.N. and Becher, P.F. ‘Sintered Al2O3/SiC–whisker composites’, Am. Ceram.Soc. Bull., 66[2] (1987) 339–342.

7. Muscat, D., Pugh, M.D., Drew, R.A.L., Pickup, H. and Steele, D. ‘Microstructure ofan extruded b-silicon nitride whisker-reinforced silicon nitride composite’, J. Am.Ceram. Soc., 75[10] (1992) 2713–2718.

8. Park, D.-S., Roh, T.-W., Han, B.-D., Kim, H.-D. and Park, C. ‘Microstructuraldevelopment of silicon nitride with aligned b-silicon nitride whiskers’, J. Euro.Ceram. Soc., 20 (2000) 2673–2677.

9. Yonezawa, T., Saitoh, S.-I., Minamizawa, M. and Matsuda, T. ‘Pressureless sinteringof silicon nitride composites’, Composites Sci. Tech., 51 (1994) 265–269.

10. Moulson, A.J. ‘Reaction bonded silicon nitride’, J. Mat. Sci., 14 (1979) 1017.11. Pugh, M.D. and Gavoret, L. ‘Nitridation of whisker reinforced reaction bonded

silicon nitride ceramics’, J. Mat. Sci., 35 (2000) 3257–3262.12. Lundberg, R.L., Kahlman, L., Pompe, R., Carlsonn, R. and Warren, R. ‘SiC-whisker

reinforced Si3N4 composites’, Am. Ceram. Soc. Bull., 66[2] (1987) 330–333.13. Dogan, C.P. and Hawk, J.A. ‘Influence of whisker reinforcement on the abrasive

wear behavior of silicon nitride- and alumina-based composites’, Wear, 203–204(1997) 267–277.

14. Campbell, G.H., Rühle, M., Dagleish, B.J. and Evans, A.G. ‘Whiskers toughening:a comparison between aluminium oxide and silicon nitride toughened with siliconcarbide’, J. Am. Ceram. Soc., 73[3] (1990) 521–530.

15. Orsini, P.G., Buri, A. and Marotta, A. J. Am. Ceram. Soc., 58 (1975) 306.16. Turner, W.E.S. J. Am. Ceram. Soc., 12 (1929) 760.17. Lee, C.J., Chae, J.I. and Kim, D.J. ‘Effect of b-Si3N4 starting powder size on elongated

grain growth in b-Si3N4 ceramics’, J. Eur. Ceram. Soc., 20 (2000) 2667–2671.18. Emoto, H. and Mitomo, M. ‘Control and characterization of abnormal growth grains

in silicon nitride ceramics’, J. Eur. Ceram. Soc., 17 (1997) 797–804.19. Sun, E.Y., Alexander, K.B., Becher, P.F. and Hwang, S.L. ‘Beta-Si3N4 whiskers

embedded in oxynitride glasses: interfacial microstructure’, J. Am. Ceram. Soc.,79[10] (1996) 2626–2632.

20. Sun, E.Y., Becher, P.F., Hsueh, C.H., Alexander, K.B., Waters, S.B., Plucknett, K.P.,Hirao, K. and Brito, M.E. ‘Microstructural design of silicon nitride with improvedfracture toughness: II, Effect of additives’, J. Am. Ceram. Soc., 81 [11] (1998) 2831–2840.

21. Jia, D.C., Zhou, Y. and Lei, T.C. ‘Thermal shock resistance of SiC whiskers reinforcedSi3N4 ceramic composites’, Ceramics International, 22 (1996) 107–112.

22. Lee, S.W., Chae, H.B., Park, D.S., Choa, Y.H., Niihara, K. and Hockey, B.J. ‘Thermalconductivity of unidirectionally oriented Si3N4/Si3N4 composites’, J. Mat. Sci., 35(2000) 4487–4493.

23. Russell, L.M., Donaldson, K.Y., Hasselman, D.P.H., Corbin, N.D. Petrovic, J.J. andRhodes, J.F. ‘Effect of vapor–liquid–solid and vapor–solid silicon carbide whiskerson the effective thermal diffusivity/conductivity of silicon nitride matrix composites’,J. Am. Ceram. Soc., 74[4] (1991) 874–877.

24. Chu, C.Y., Singh, J.P. and Routbort, J.L. ‘High-temperature failure mechanisms ofhot-pressed Si3N4 and Si3N4/Si3N4-whisker-reinforced composites’, J. Am. Ceram.Soc., 76[5] (1993) 1349–1353.


25. Chu, C.Y. and Singh, J.P. ‘Mechanical properties and microstructure of Si3N4-whisker-reinforced Si3N4 matrix composites’, Ceram. Eng. Sci. Proc., 11[7–8] (1990) 709–720.

26. Becher, P.F., Sun, E.Y., Plucknett, K.P., Alexander, K.B., Hsueh, C.H., Lin, H.T.,Waters, S.B., Westmoreland, C.G., Kang, E.S., Hirao, K. and Brito, M.E.‘Microstructural design of silicon nitride with improved fracture toughness: I, Effectsof grain shape and size’, J. Am. Ceram. Soc., 81[11] (1998) 2821–2830.

27. Shih, C.J., Yang, J.-M. and Ezis, A. ‘Microstructure and properties of reaction-bonded/hot-pressed SiCw/Si3N4 composites’, Composites Sci. Techn., 43 (1992)13–23.

28. Bellosi, A. and De Portu, G. ‘Hot-pressed Si3N4-SiC whisker composites’, Mat. Sci.Eng., A109 (1989) 357–362.

29. Pezzotti, G., Tanaka, I. and Okamoto, T. ‘Si3N4/SiC-whisker composites withoutsintering aids: III, High temperature behaviour’, J. Am. Ceram. Soc., 74[2] (1991)326–332.

30. Becher, P.F. ‘Microstructural design of toughened ceramics’, J. Am. Ceram. Soc., 74(1991) 255.

31. Faber, K.T. and Evans, G.A. ‘Crack deflection processes: theory and experiment’,Acta Metall., 31[4] (1983) 565–584.

32. Peillon, F.C. and Thevenot, F. ‘Microstructural designing of silicon nitride related totoughness’, J. Eur. Ceram. Soc., 22 (2002) 271–278.

33. Baldacim, S.A., Santos, C. Silva, O.M.M. and Silva, C.R.M. ‘Ceramics compositesSi3N4-SiC(w) containing rare earth concentrate (CRE) as sintering aids’, Mat. Sci.and Eng., A367 (2004) 312–316.

34. Rajan, K. and Šajgalík, P. ‘Microstructurally induced internal stresses in b-Si3N4

whiskers-reinforced Si3N4 ceramics’, J. Eur. Ceram. Soc., 17 (1997) 1093–1097.35. Baldacim, S.A., Santos, C., Silva, O.M.M. and Silva, C.R.M. ‘Mechanical properties

evaluation of hot-pressed Si3N4-SiC(w) composites’, Int. J. Refrac. Met. and HardMat., 21 (2003) 233–239.

36. Nixon, R.D., Chevacharoenkul, S., Davis, R.F., Huckabee, M.L. and Buljan, S.T.‘Deformation behaviour of SiC whisker reinforced Si3N4’, in Materials ResearchSociety Symposium Proceedings, vol. 78, ed. Becher, P.F., Swain, M.V. and Somia,S. Materials Research Society, Pittsburg, PA (1987) 295–302.

37. Backhaus-Ricoult, M., Castaing, J. and Roubort, J.L. ‘Creep of SiC-whiskers reinforcedSi3N4’, Revue Phys. Appl., 23[3] (1988) 239–249.

38. Nixon, R.D., Koester, D.A., Chevacharoenkul, S. and Davis, R.F. ‘Steady-state creepof hot-pressed SiC whisker-reinforced silicon nitride’, Composites Sci. Tech., 37(1990) 313–328.

39. Desmarres, J.M., Goursat, P., Besson, J.L. Lespade, P. and Capdepuy, B. ‘SiC whiskersreinforced Si3N4 matrix composites: Oxidation behavior and mechanical properties’,J. Eur. Ceram. Soc., 7 (1991) 101–108.

40. Zhu, S., Mizuno, M., Kagawa, Y., Nagano, Y. and Kaya, H. ‘Static and cyclic fatiguein SiC whisker-reinforced silicon nitride composite’, Mat. Sci. and Eng., A251(1998) 113–120.

41. Kaya, H. ‘The application of ceramic-matrix composites to the automotive ceramicgas turbine”, Composites Sci. and Tech., 59 (1999) 861–872.

42. Bhola, R., Das Gupta, S. and Jacobs, J.K. ‘Ceramic cutting tool inserts’, KeyEngineering Materials, 122–124 (1996) 235–246.

43. Brandt, G. ‘Advanced tool materials’, Euro PM 2001 Proceeding, 1 (2001) 90–95.

58

3.1 Introduction

The traditional or conventional ceramics are generally in monolithic form.These include bricks, pottery, tiles and a variety of art objects. The advancedor high-performance monolithic ceramic materials represent a new andimproved class of ceramic materials where, frequently, some sophisticatedchemical processing route is used to obtain them. Generally, their characteristicsare based on the high quality and purity of the raw materials used. Examplesof these high-performance ceramics include oxides, nitrides, carbides ofsilicon, aluminium, titanium and zirconium, alumina, etc.

Monolithic high-performance ceramics combine some very desirablecharacteristics, e.g. high strength and hardness, excellent high-temperaturecapability, chemical inertness, wear resistance and low density. They are,however, not very good under tensile and impact loading, and, unlike metals,they do not show any plasticity and are prone to catastrophic failure undermechanical or thermal loading (thermal shock). The difference in the behaviourof metals and ceramics can be categorized by saying that metals are forgivingwhile ceramics are not forgiving. The forgiving nature of metals has itssource in the high mobility of dislocations of their atoms in them, whichallows them to deform plastically before fracture. Plastic deformation beingan energy-absorbing process, the fracture process in metals involves extensiveenergy dissipation. Ceramic materials are lacking such energy-dissipatingphenomenon, which causes them to fail in a catastrophic fashion, i.e. makesthem unforgiving. Therefore, a critical need exists for increasing the toughnessof ceramic materials. With a view to achieving high fracture toughness ofceramic materials, the major effort of the materials community in the field ofstructural materials has been directed towards incorporating a variety ofenergy-dissipating phenomena in the fracture process of ceramics, i.e. impartingdamage-tolerant behaviour. Improving the toughness and in-service reliabilityof ceramic materials is thus the major objective. One of the importantapproaches to attaining these goals is through composite and glass/glass-

3Fibre-reinforced glass/glass-ceramic

matrix composites

R B A N E R J E E and N R B O S E, Central Glass andCeramic Research Institute, India

Fibre-reinforced glass/glass-ceramic matrix composites 59

ceramics matrix composites, and these are among the materials attractingmajor attention today.

During the last quarter of the twentieth century, materials scientists havebeen able to revolutionize a major change in the development of materialsthat are considered suitable for applications at temperatures of more than300oC with characteristics of high fracture toughness. In the past, designersworked with data on the properties of homogeneous, isotropic materials anddesigned their components to fit the ranges of ‘design allowances’ by usingfibre-reinforced polymer matrix and metal–matrix composites for structuralapplications. The use of such composites increased to the point where boron-and graphite-fibre reinforced epoxy resin and boron-reinforced aluminiumwere restricted to applications up to 300∞C. Recently, however, the conceptof composite materials has permitted almost limitless tailoring of compositesto create entirely new designs never previously possible.

By selecting judiciously the types of material constituents, their relativepercentages, their orientation and fabrication methods, the designer can nowwork closely with the materials scientist to optimize system performance.The widespread acceptance of this philosophy, combined with new challengesto create engineering advances in aerospace and commercial areas, has createdtremendous opportunities for new composites possessing greater environmentalstability, higher temperature capabilities and high fracture toughnesscharacteristics. The quest for such excellent environmental aspects andimproved characteristics of materials prompted the development of fibre-reinforced glass/glass-ceramic matrix composite (CMC). The combinationof starting materials with suitable properties and appropriate fabricationprocedure ultimately determines the properties of the resultant composite.Although this simple and obvious statement encompasses too many facets tobe considered to achieve the high-performance characteristics of the resultantcomposite, two facets can be considered for comment. First, the fibres shouldnot be greatly degraded during processing either by handling or by chemicalreaction, and second, the resultant fibre–matrix interface must have thecharacteristics necessary to prevent excess fibre–matrix bonding.

As should be clear from the discussion above, high-performance ceramicsmust have superior structural and/or mechanical characteristics because theyfind application in some very demanding environments, e.g. rocket nozzles,heat exchangers, automobile engines and cutting tools. Yet another importantfactor is the cost of ceramics. The great challenge is to produce consistentand reliable ceramic components having superior properties but withoutexcessive cost, i.e. they should be competitive on a cost/performance basiswith the materials they seek to replace. In this regard fibre-reinforced glass-ceramic composite has taken a leading role.

Incorporation of fibres, whiskers or particles in a ceramic/glass-ceramicmatrix can result in a tough ceramic material. This happens because


incorporation of reinforcements introduces energy-dissipating phenomenasuch as debonding at the fibre/matrix interface, crack deflection, fibre bridging,fibre pullout, etc. In this regard, proper control of the characteristics of theinterface region is of obvious importance. Yet another point to note is thatwhile the ratio of the modulus of the reinforcement and the polymer or metalmatrix is generally between 10 and 100, this ratio for a CMC is rather low,and can frequently be equal to unity or even less. The fibre may be continuousor discontinuous with a high aspect ratio. The high modulus ratio in polymer–matrix composites (PMC) and metal–matrix composites (MMC) allows anefficient load transfer from the matrix to the fibre via a strong interface.However, in a CMC, unlike PMC and MMC, the low modulus ratio meansthat the reinforcement and the matrix are not very different in their load-bearing capacity, i.e. a simple increase in strength of a ceramic material israrely the objective. It is therefore necessary to consider low modulus basedmatrix material (e.g. glass/glass-ceramic matrix) in CMC for developinghigh strength in the materials.

3.2 Types of fibre suitable as reinforcements in

different glass/glass-ceramic matrix composites

Reinforcements in the form of continuous fibres, short fibres, whiskers orparticles are available commercially. Continuous ceramic fibres are veryattractive as reinforcements in high-temperature structural materials. Theyprovide high strength and elastic modulus with high temperature-resistantcapability and are free from environmental attack. Ceramic reinforcementmaterials are divided into oxide and non-oxide categories, listed in Table3.1. The chemical compositions of some commercially available oxide andnon-oxide reinforcements are given in Table 3.2 and Table 3.3.

Table 3.1 Ceramic reinforcement materials56

Category Reinforcement materials

Continuous fibresOxide Al2O3, (Al2O3 + SiO2), ZrO2

Silica based glasses, etc.Non-oxide B, C, SiC, Si3N4, BN

Discontinuous fibresWhiskers SiC, TiB2, Al2O3

Short fibres Glass, Al2O3, SiC, (Al2O3 + SiO2)Vapour-grown carbon fibre

Particles SiC, TiC, Al2O3


Many researchers have used glass and glass-ceramic matrices for reinforcingwith high-modulus graphite fibres [1, 2], silicon carbide fibres and siliconcarbide mono-filaments [3–7]. Very strong, tough and refractory compositeswere obtained from these efforts.

Incorporation of discontinuous graphite fibres in glass and glass-ceramicmatrices has also been reported [8, 9]. The basic mechanism of strengtheningin fibre composite is that of load transfer by the matrix to the fibres throughshear. This load transfer takes place at the fibre ends within a few fibrediameters. Depending on the length of the fibre, the amount of load transferredby the matrix to the fibre changes. Therefore, the fibre may not be loaded tothe breaking stress and full advantage cannot be taken of its reinforcingability. The stress concentration at the fibre ends also results in lower strengththan would be possible. The load transfer depends on the properties of the

Table 3.2 Composition of oxide ceramic reinforcements56

Trade name Manufacturer Composition (%)

Al2O3 ZrO2 SiO2 B2O3 Fe2O3

ContinuousNextel 312 3M 62 — 24 14 —Nextel 440 3M 70 — 28 2 —Nextel 480 3M 70 — 28 2 —Nextel 550 3M 73 — 27 — —Nextel 610 3M >99 — 0.2–0.3 — 0.4–0.7Astroquartz J.P. Stevens — — 99.95 — —Saphikon Saphikon Al2O3 single crystal

Discontinuous (various silica-based glasses)Saffil ICI 96 — 4 — —Fibermax Carborundum 72 — 28 — —Fiberfrax Carborundum 52 — 48 — —

Table 3.3 Composition of non-oxide ceramic reinforcements56

Trade name Manufacturer Composition (%)

Continuous fibresNicalon Nippon Carbon Co. SiC + O + CSCS, Sigma Textron Specialty SiC on tungsten

Materials, BP or carbon substrateTyranno Ube Ind. SiC + Ti + CTyranno Ube Ind. SiC + Zr + C

Discontinuous fibresSilar Adv. Composite SiC whiskers

Material Corp.Tokawhisker Tokai Carbon SiC whiskers


matrix, the reinforcing fibre and the interface between them. These propertieswill determine the extent of reinforcement that can be put into the compositewith the given aspect ratio of the fibre [10]. The salient features of oxidefibres and non-oxide fibres and their fabrication process are discussed below.

3.2.1 Oxide fibres

Ceramic oxide fibres, both continuous and discontinuous, have beencommercially available since the 1970s, and processing and microstructurecontrol are very important in obtaining the desired properties. Among thedesirable characteristics in any ceramic fibre for structural applications are:

∑ High theoretical density, i.e. low porosity∑ Small grain size for low-temperature applications∑ Large grain size for high-temperature applications∑ High purity.

Alumina fibres have g, d, h and a allotropic forms. a-Alumina is thethermodynamically stable form. In practice, it is very difficult to control thetime and temperature conditions to proceed from g to a. At low firing, thefibres will give a smaller grain size and therefore an unacceptable level ofporosity. At higher processing temperatures, porosity can be eliminated butexcessive grain growth will result. This dilemma can be avoided by introducinga second phase that restricts grain boundary mobility while the porosity isremoved at high temperature. It is possible to select the type and amount ofthe second phase that inhibits the grain growth at the service temperature.There are various ways to select this second phase by trial and error. It hasbeen made possible to lower the working temperature by introducing oxidesof silicon, phosphorus, boron or zirconium as the second phase, therebyinhibiting the formation of a grain boundary. Some results obtained afterexperimental studies are a-alumina + 15–20% ZrO2, d-alumina + 4% SiO2

and a-alumina + 0.4% Fe2O3 + 0.25% SiO2 for lowering the workingtemperature during the formation of fibres.

Various types of oxide fibres available commercially can be considered assuitable reinforcements other than the types of fibres listed in Table 3.2:

∑ Fibre FP: a polycrystalline continuous a-alumina fibre yarn produced byDuPont in the 1980s.

∑ Fibre PRD-166: another polycrystalline continuous alumina fibre yarnproduced by DuPont in the 1980s. PRD-166 fibre yarn is a modified formof FP fibre yarn. The diameter of this fibre filament is 20 mm. Themodification of FP fibre is made by incorporating 15–20 wt% yttria-stabilized zirconium particles. The idea of incorporating Y2O3-stabilizedzirconia particles was to take care of problems such as unstable mechanical


properties caused by grain growth and creep at high temperature. PRD-166 fibre had a rough surface and an average grain size of about 0.5 mm.The zirconia particles were about 0.1 mm across and located mostly atgrain boundary triple points. Their function was to inhibit grain growth inalumina fibre. Although DuPont no longer produces these fibres, theirfabrication represented an important step in the processing of alumina-type fibres [11–13].

∑ Minnesota Mining and Manufacturing Co., also known as 3M Co., hasdeveloped an a-alumina fibre, trade name Nextel 610, via the sol-gelroute. Figure 3.1 shows the 3M process schematically.

∑ Sumitomo Chemical Co. produces a fibre that is a mixture of alumina(85%) and silica (15%). The fibre structure consists of fine crystallites ofspinel. SiO2 serves to stabilize the spinel structure and prevents it fromtransforming to a-alumina [14]. The flow diagram of this process is shownin Fig. 3.2.

∑ Series of various Nextel fibres produced by 3M Co. are mainly Al2O3 +SiO2 and some B2O3. The properties of some oxide fibres and monoxidefibres are given in Table 3.4.

ICI Co. uses a sol-gel method to produce silica-stabilized alumina (Saffil)and calcia-stabilized zirconia fibre [15]. The saffil fibre is a d-alumina shortstaple fibre that has about 4% SiO2 and a very fine diameter (3 mm).

Filter

PumpSpinneret

Reservoir oforganic basicAl salt solution

Drawingwheels WinderPyrolysis

furnace zone(1400∞C)

Low temperaturefurnace zone forstraightening fibre

3.1 3M’s process for making Al2O3 fibre (reproduced by permission ofChapman & Hall)56.


Sowman has provided details of the process used by 3M Co. for makingthe Nextel oxide fibre [16]. The starting material for this fibre is aluminium

acetate, Al(OH)2OOCCH2 . 13

H3BO3. It is a product of Niacet Corporation

under the trade name Niaproof.A continuous polycrystalline a-alumina, trade name Almax, has been

prepared by researchers at the Mitsui Mining Co. [17] by dry spinning aviscous slurry consisting of an aluminium salt, a fine powder of intermediatealumina, and an organic binder to produce the precursor fibre; this is followedby prefiring (calcining) and firing (sintering) the precursor fibre to producean alumina fibre. Figure 3.3 shows the flow diagram of Almax alumina fibre.

A continuous monocrystalline sapphire (Al2O3) fibre has been preparedas single-crystal fibres by LaBelle and Mlavsky using a modified czochralskipuller and radio frequency heating. The technique adopted in this method iscalled edge-defined film-fed growth (EFG) [18–22]. Figure 3.4 shows aschematic of the EFG method.

Polyaluminoxanes–Sicompound

Alkoxy alumino gelcompound AIR3

Al O

R

PolymerizationreactorAIR3 + H2O

Alkyl silicateand organic

solvent

Dryspinningzone

Winder for drawingin organic fibre

Al2O3: 70–100%SiO2: 30–0%

3.2 Flow diagram of the Sumitomo process for making a mixture ofalumina and silica fibre (reproduced by permission of Chapman &Hall)56.

Fibre-reinforced glass/glass-ceramic m

atrix composites

65

Table 3.4 Composition and properties of some oxide fibres56

Fibre type Composition (wt%) Diameter Density Tensile strength Tensile modulus——————————————————————— (mm) (gm/cm3) (MPa) (GPa)Al2O3 SiO2 B2O3 Fe2O3

3M Nextel 312 62 24 14 — 10–12 2.7 1700 1523M Nextel 440 70 28 2 — 10–12 3.05 2000 1863M Nextel 480 70 28 2 — 10–12 3.05 2070 2203M Nextel 550 73 27 — — 10–12 3.03 2240 2203M Nextel 610 79 0.2–0.3 — 0.7 10–12 3.75 1900 370

ICI Saffil 96 4 — — 3 2.3 1000 100Saphikon Al2O3 single crystal 70–250 3.8 3100 380Sumitomo 85 15 9 3.2 2600 250


Several investigators using a laser-heated floating zone method haveprepared a variety of ceramic fibres. Gasson and Cockayne [22] used laserheating for the crystal growth of Al2O3, Y2O3, MgAl2O4 and Na2O3. Haggerty[23] used a four-beam heated float zone method to grow single-crystal fibresof Al2O3, Y2O3, TiC and TiB2. The laser-heated float zone technique isshown in Figure 3.5.

1500–3000 poise at 25∞C

50–100 m/min

250–500∞C

1400–1600∞C

Al2(OH)5Cl Al2(OH)5Cl powder

AlCl36H2O

Spinning andsintering aid

Dispersion

Mixing

Filtration

Forming gel

Spinnablemixture

Dry spinning

Precursorfibre

Prefiring

Firing

a-Alumina fibre

3.3 Flow diagram for Almax Alumina fibre (reproduced bypermission of Chapman & Hall)56.


Another novel technique of making oxide fibres is called the inviscid melttechnique [24]. In principle, any material in a molten state can be drawn intoa fibre shape. Organic polymeric fibres such as nylon, aramid, etc., as wellas a variety of glasses can be routinely converted into fibrous form bypassing a molten material, having an appropriate viscosity, through an orifice.The inviscid (meaning low viscosity) melt technique uses this principle tomake oxide fibres.

Moltenalumina

Sapphireseed

Molybdenumcapillary

3.4 Edge-defined, film-fed, growth process for making a single-crystalalumina fibre (reproduced by permission of Chapman & Hall)56.

Pull

Feedsource

rod

Seed crystal

Laser

Molten

3.5 Laser-heated float zone technique (reproduced by permission ofChapman & Hall)56.


3.2.2 Non-oxide fibres

Commercially available non-oxide ceramic reinforcements are in threecategories: continuous, discontinuous, and whiskers. The great breakthroughin the ceramic fibre area has been the concept of pyrolysing polymers undercontrolled conditions, containing the desired species to produce high-temperature ceramic fibres. Silicon carbide fibre is a major development inthe field of ceramic reinforcements.

Non-oxide fibres via polymers

The SiC fibre obtained via CVD is very thick and not very flexible. By analternative route, very fine, continuous and flexible fibre was obtained byYajima and his colleagues [25, 26] in Japan using a process of controlledpyrolysis of polymeric precursor. The ceramic fibres produced by this processpossess good mechanical properties; good thermal stability and oxidationresistance have enormous potential for the development of ceramic matrixcomposites. Figure 3.6 shows the various steps involved in processing non-oxide fibres through the polymeric route, and Fig. 3.7 shows schematicallythe Yajima process of making SiC fibre from a polycarbosilane.

Non-oxide fibres via CVD

Silicon carbide (SiC) deposited on a substrate of tungsten or carbon heatedto about 1300oC [27] is called sigma fibre (BP Sigma). A detailed schematicof the process used by BP to make its Sigma fibre is shown in Fig. 3.8.Textron Specialty Material Co. has developed a series of surface-modified

Curing

Melt orsolutionspinning

Controlledpyrolysis

Polymer precursor

Precursor fibre

Cured or stabilized fibre

Ceramic or glass fibre

3.6 Various steps involved in processing non-oxide fibres through thepolymeric route (reproduced by permission of Chapman & Hall)56.


silicon carbide fibres, called SCS fibres. These fibres have a complex through-thickness gradient structure. SCS fibre is a thick fibre (142 mm) and isproduced by CVD of silicon- or carbon-containing compounds onto a pyrolyticgraphite-coated carbon core. The pyrolitic graphite coating is applied to acarbon monofilament to give a substrate of 37 mm. SiC is then coated byCVD to give a final monofilament of 142 mm diameter. Figure 3.9 showsschematically the cross-sections of the two Textron SCS-type SiC fibres.

Dichlorodimethylsilane

Si

Cl

Cl

CH3

CH3

Polydimethylsilane

CH3

Si

CH3 n

Polycarbosilane

CH3

Si

CH3

C

H

H

n

Dechlorinationwith Na (to NaCl)

Polymerization at470∞C in autoclave

Melt spinning at350∞C (N2)

Curing at 190∞C inair or RT in ozone

Pyrolysis to 1300∞Cin vacuum (1000∞C h–1)

Polycarbosilanefibre

SiC fibreAmorphous or micro-

crystalline b-SiC

Polycarbosilane fibreswith molecular cross-linking by oxygen to

avoid subsequentmelting

3.7 Yajima process for making SiC from a polycarbosilane(reproduced by permission of Chapman & Hall)56.


W. filament

SiC/W

+

+

Exhaust

Scrubber

Refrigeration

Distillation

Waste by-product

Gases forrecuperation

Hydrogensupply

Flowmeters

Silanevaporizer

Silanesupply

Photooptical

diametersensor

3.8 Schematic process for making SiC monofilament fibre by the CVDmethod (reproduced by permission of Chapman & Hall)56.

Pyralytic graphite-coated carbon core

Inner zonecarbon-richb-SiC

Outer zone:stoichiometricb-SiC

Carbon-richsurface coating(0–4 mm)

SCS-6 (~140 mm) SCS-0 (~75 mm)

3.9 Schematic of two Textron SCS-type silicon carbide fibres(reproduced by permission of Chapman & Hall)56.


Nicalon-type SiC fibres

The Nicalon fibre (10–20 mm) available commercially consists of a mixtureof b-SiC free carbon and SiO2 [28]. The properties of Nicalon start to degradeabove about 600oC because of the thermodynamic instability of the compositionand microstructure. Ceramic-grade Nicalon fibres, designated the NL series,having low oxygen content are also available.

Other SiC/Si3N4-type fibres

Various SiC-type fibres with elemental compositions of Si–C, Si–N–C–O,Si–B-N, Si–C–O and Si–Ti–C–O are commercially available. These fibresare made from polymeric precursors. A multifilament SiC fibre, called Tyranno,is produced by Ube Industries, Japan [29]. This fibre is made by the pyrolysisof poly(titano carbosilanes) and contains 1.5–4% titanium by weight. Anothermultifilament fibre is called silicon carbonitride, trade name HPZ, producedby Dow Corning Corporation, USA.

Several researchers also produce silicon nitride fibre by using variousorganometallic compounds. SiC-based silicon nitride fibre is produced bythe pyrolysis of organosilazane polymers with methyl groups on Si and N[30]. Researchers at Tonen Corporation of Japan also produce pure siliconnitride fibre. This fibre is made by using perhydropolysilazane polymer(PHS). Another silicon nitride fibre, called Tonen SiNB, is based on boron.Silicon nitride (Si3N4) fibres can also be prepared by reactive chemicalvapour deposition (CVD) using volatile silicon compounds. The reactantsare generally SiCl4 and NH3. Si3N4 is deposited on carbon or tungsten substrate.All these fibres can be used for the fabrication of glass/glass-ceramic matrixcomposite with proper interface modification. The interface modification isoutside the scope of this discussion.

Whiskers

Whiskers are normally obtained by vapour phase growth. They aremonocrystalline, short fibres with extremely high strength because of theirhigh aspect ratio (50 to 10 000). They have a diameter of a few microns, butthey do not have uniform dimensions and properties.

Other non-oxide fibres

There are other promising ceramic fibres, e.g. boron carbide and boronnitride. Boron nitride fibre has the same density (2.2 g cm–3) as carbon fibre,but has a greater oxidation resistance and excellent dielectric properties.Boron carbide fibre is a very light and strong material.


Production

SiC particles and whiskers are produced from rice hulls. The rice hulls areheated at about 700∞C in the absence of oxygen. This system is called coking.The coked rice hulls consist of equal amounts of C and SiO2. When thecoked rice hulls are heated at a temperature of 1500–1600∞C for about anhour in an inert atmosphere (N2 or NH3 gas) they form SiC whiskers. Figure3.10 shows a schematic of the manufacturing process for SiC whiskers bythe Vapour–Liquid–Solid (VLS) processing method. In this process siliconand carbon are obtained from SiO and CH4 gases respectively. Figure 3.11shows the chemical transformation of the VLS whisker process.

3.3 Methods for manufacturing different fibre-

reinforced glass/glass-ceramic matrix

composites

The techniques for fabricating glass/glass-ceramic matrix composites arebased on the final properties and performance of the resultant composite.

Dispersion ofshreddedrice hull

Carbon tubereactor

Rice hull

Whisker andcarbon

separation

Whisker andhull relictseparation

Drying

Carbonoxidation

SiC whisker

3.10 Manufacturing process for silicon carbide whiskers by VLS fromrice husk (reproduced by permission of Chapman & Hall)56.


Some of the techniques are unconventional and few are conventional powderprocessing techniques. The techniques are described below.

3.3.1 Conventional techniques

Cold pressing and sintering

The first stage in producing a component is to press the powder and thebinder into a desired shape using a sufficiently high pressure so that a relativelydense and strong green body is formed. In certain circumstances, a variety offast production methods can be used such as extrusion, blow moulding,injection moulding, etc. It is necessary to remove the organic binder beforea fully sintered body with a near-theoretical density can be obtained. Thegreatest uniformity of density is obtained by the application of pressure fromall directions, which is known as isostatic pressing. To improve the strengthof the resultant composite, compacted ingredients have to be heated to elevatedtemperatures in order to burn off the binder and to consolidate the powderfurther by sintering. Generally, in the sintering step, the matrix shrinksconsiderably and the resultant composite develops many cracks. There areother limitations on sintering of glass, ceramic or glass-ceramic matrix materialscontaining high aspect ratio reinforcements. Because of the difference in

SiC

Whisker

Si + C = SiC

CatalystCO(g)

SiO(g) + C = Si + CO(g)

Fe–Si–C

2H2(g)

SiO(g)

Co(g)

SiO2and C

Generator

CH4(g)

CH4(g) = C + 2H2(g)

3.11 Chemical transformation of the vapour-liquid-solid (VLS) phasesfor growing whiskers (reproduced by permission of Chapman &Hall)56.


thermal expansion coefficients of the reinforcement and matrix, a hydrostatictensile stress may develop in the matrix on cooling. This effect retards thesintering process [31, 32]. In glass matrix composites sintering is retarded ifthe reinforcement is greater than about 15 vol% [33, 34]. Yet another limitationto consider is the bridging phenomenon, which is caused by the formation ofnetworks resulting from whiskers or fibres. So it is important to optimize thevolume fraction as well as the aspect ratio of the whiskers or reinforcementfor getting the densified product [35]. Experimental evidence was establishedby the fact that the sintered density of silicon carbide reinforced aluminadecreased as the aspect ratio of the whisker increased [36].

Hot-pressing

Simultaneous application of pressure and high temperature can acceleratethe rate of densification and production of a pore-fee and fine-grained compactof the ceramic materials. The most important technique used to producecontinuous fibre-reinforced glass and glass-ceramic composite is the slurryinfiltration process [37–40]. This was developed about 20 years ago in theUnited Kingdom for the production of glass-matrix composites but it is nowalso widely used for glass–ceramic matrix composites. The intimate mixingof continuous fibres and the glass or glass–ceramic matrix is achieved bydrawing bundles of fibres, called tows, through a slurry of powdered glass inwater and a water-soluble resin binder. Figure 3.12 shows a schematic of this

Glass slurrytankFibres

Glass-impregnatedfibre tape

Stack of glass-impregnatedfibre tapes

Binder burnout500∞C

Graphite dieHot pressing

800–925∞C

Pressure

Fibre/glasscomposite

3.12 Schematic of the slurry infiltration process for making a fibre-reinforced glass and glass–ceramic composite (reproduced bypermission of Woodhead Publishing Limited)74.


process. The tows, impregnated with the slurry, are wound onto a mandrel toform a monolayer tape. The tape is cut into plies, which are stacked into therequired stacking sequence before burnout of the binder. Wetting agents maybe added to ease the infiltration of the fibre tows. The impregnated tape isthen hot-pressed for consolidating the matrix. The CMCs produced by hotpressing are of a superior quality. The flaws encountered in CMCs producedby hot-pressing may be due to the presence of a combination of matrix-richregions and fibre-rich regions. This kind of inhomogeneity weakens thecomposite. Some CMCs produced by various investigators using hot-pressingtechniques are silicon carbides, alumina and carbon fibres in a variety ofglass, glass–ceramic and oxide–ceramic matrices [39–41]. The slurry infiltrationprocess is ideal for making glass or glass–ceramic matrix composites becausethe processing temperatures for these materials are lower than those used forcrystalline matrix materials. Figure 3.13(a) shows an optical micrograph ofa transverse section of a unidirectional alumina fibre/glass matrix composite,while Fig. 3.13(b) shows the pressure and temperature schedule usedduring hot-pressing of this composite. Some porosity can be seen in thispicture.

Lanxide process

The Lanxide process was developed by Lanxide Corporation [42]. It involvesthe formation of a ceramic matrix by the reaction between a molten metaland a gas. For example, when molten aluminium is reacted with oxygen,alumina is formed. The ceramic matrix occurs outwards from the originalmetal surface. In the case of fibrous composites, filament winding or a fabriclay-up may be used to make a preform. A fabric made of a continuous fibrecan also be used. The fabric is coated with a proprietary chemical solution toprotect the fibre from highly reducing aluminium and to provide a weakinterface. A barrier is placed on the preform surfaces to control the rate ofgrowth of the matrix material. In this method, the transport of the ceramicmatrix occurs continuously at the oxidation reaction front. The desired reactionproduct forms on the surface of the molten metal and grows outward. Thisprocess is considered as a low-cost process and there is a possibility to getnear-net shape products. It has been reported [42] that products with goodmechanical properties (strength, toughness, etc.) can be obtained by thisprocess, which can make glass/glass–ceramic matrix composites. A schematicof the Lanxide process is shown in Fig. 3.14.

Various researchers have used different techniques for the fabrication ofceramic matrix composite (CMC) by using (a) chemical vapour deposition(CVD), (b) chemical vapour infiltration (CVI) and (c) modelling of CVI[43–54], but it is difficult to fabricate glass/glass–ceramic composite usingthese techniques.


Sol-gel technique

One of the most innovative approaches to ceramic and glass processing isthe sol-gel technique. A brief description of the process is given below. Thesol-gel route of making any glass or ceramic involves the formation of the

(a)

TemperaturePressure

0 50 100 150 200Time (min)

(b)

1000

800

600

400

200

0

Tem

per

atu

re (∞C

)

6

5

4

3

2

1

0

Pre

ssu

re (

MP

a)

3.13 (a) Optical micrograph of a cross-section of a unidirectionalalumina fibre/glass matrix composite made by slurry infiltration; (b)pressure and temperature schedule used during hot-pressing of thiscomposite (reproduced by permission of Chapman & Hall)56.


appropriate glass or ceramic structure by chemical polymerization of suitablecompounds in the liquid state (sol) at low temperatures, followed by chemicalreactions such as hydrolysis or condensation at temperatures much lowerthan those used in powder processing or direct melting. The particle size insol generally varies between 1 and 100 nm. It can also be obtained by mixinga metal-containing precursor (e.g. an acid) and water. Hydrolysis and pre-condensation reactions make the sol viscosity increase until a gelled state isobtained. This gel is like a wet solid and is termed a precursor. The wet gelis dried to remove any unwanted residue (water, organic compounds, etc.).The gel is then converted into glass or ceramic by heating at temperaturesmuch lower than those used in direct melting processes. The glass or ceramicthus obtained may be in the form of powder, film, fibre, etc. The slurryinfiltration process results in a fairly uniform fibre distribution inside thecomposite and low porosity, and enhanced high-strength values of thecomposites. Figure 3.15 shows the sol-gel process flow diagram.

For the fabrication of ceramic matrix composite conventional polymerhandling and processing techniques are used. Fibrous preforms are made byimpregnation of sol in vacuum or filament winding techniques. In filamentwinding, fibre tows or ravings are passed through a tank containing sol andthe impregnated tows are wound on a mandrel to a desired shape and thickness.The sol is converted to gel and the structure is removed from the mandrel. Afinal heat treatment is done at 1400oC to get ceramic or glass–ceramic

(a)

Molten metal

(b)

gSL

gSVAir

Infiltratedpreform

Reactive gaseousenvironment

Preform

Liquid metal

3.14 Lanxide process: (a) infiltration of preform; (b) wicking of liquidmetal along grain boundaries (reproduced by permission ofWoodhead Publishing Limited)74.


composite. A fibrous preform can be vacuum impregnated with a polymerprecursor in liquid form and an appropriate heat treatment then given toconvert the polymer into ceramic. Particulate or whisker reinforcement canbe disposed in the sol or gel state. Sol-gel filament winding of a fibre preformand vacuum impregnation of woven preforms are shown in Fig. 3.16 andFig. 3.17 respectively.

Researchers at GEC in England have emphasized the fabrication ofcontinuous fibre-reinforced ceramics in various complex shapes. They haveused liquid precursors to produce a ceramic matrix in a fibrous preform [55,56].

Sol-gel processing is a viable means of preparing glass, ceramics and thinfilms through hydrolysis and condensation of metal alkoxides in organicsolvents [57–61]. Compared with conventional techniques, the sol-gel methodhas several advantages because many multi-component oxides can be preparedwith a higher degree of chemical purity and easier control of stoichiometry.

The sol-gel technique offers excellent composition control, low-temperatureprocessing and short fabrication times at comparatively low cost. One of itsdisadvantages consists of the fact that only a small thickness (up to ~ 200nm) of high-quality film can be achieved in one coating cycle, so several

MixMix sol with reinforcementor gel with reinforcement

PourPour sol over

preform

Dry Dry

Hot press

CalcineHeat to producerequired ceramic

RepeatRepeat infiltration and drying

until required density

Fire Film

Fibre

Powder

(a) (b)

3.15 Sol-gel processing: (a) infiltration of a preform; (b) mixingreinforcement in a sol or a gel (reproduced by permission ofWoodhead Publishing Limited)74.


Dryer

Sol

Spool Filamentwinding

Gelled body

Heating coils

Heat to convert the gelinto glass or ceramic

3.16 Sol-gel filament winding of a woven preform (reproduced bypermission of Chapman & Hall)56.

Vacuum

Sol

Preform

Vacuum impregnation by sol

Gelation

Heatingcoils

Heat to convert gel into glass or ceramic

3.17 Vacuum impregnation of a woven preform (reproduced bypermission of Chapman & Hall)56.


coating cycles are necessary when thick films are required. One of the waysto make a thick layer is to form a composite material. Several researchershave reported the preparation of cermets [62] or sol-gel polymeric composite[63, 64]; lead zirconate titanate (LZT) and ZrO2 thick films have beensynthesized [57] using the dispersion of ceramic powders in sol containingzirconium alkoxides.

3.4 Properties of glass/glass–ceramic matrix

composites

The important ceramic matrix materials are glass, silicon carbide, siliconnitride, alumina, glass–ceramics, sialons, intermetallics and some elementalmaterials. A list of some ceramic matrix materials is given in Table 3.5.

The characteristic high strength and brittleness of ceramic matrix materialscan be judged by the types of bonding in their structure [65, 66]. In ceramicmatrix materials with ionic bonding, there occurs a transfer of electronsbetween the atoms, and in case of covalent bonding, the electrons are sharedbetween atoms. The properties of some ceramic matrix materials are givenin Table 3.6.

Table 3.5 List of some ceramic matrix materials56

Nitrides Silicon nitride (Si3N4), boron nitride (BN)

Carbides Silicon carbide (SiC), boron carbide (B4C), titanium carbide (TiC)

Mixed oxides Mullite (3Al2O3.2SiO2), spinel (MgO.Al2O3)

Single oxides Alumina (Al2O3), zirconia (ZrO2), titania (TiO2), magnesium oxide(MgO), Silica (SiO2)

Intermetallics Nickel aluminide (NiAl, Ni3Al), titanium aluminide (TiAl, Ti3Al),molybdenum disilicide (MoSi2)

Elements Carbon (C), boron (B)

Table 3.6 Properties of some ceramic matrix materials56

Ceramic matrix Physical and mechanical propertiesmaterials ——————————————————————————————

Density Melting Young’s Coefficient of Fracturer (g.cm–3) point modulus thermal toughness

(∞C) E (GPa) expansion KIC

a (10–6K–1) (MPa. m1/2)

Al2O3 3.9 2050 380 7–8 1–3SiC 3.2 — 420 4.5 2.2–3.4Si3N4 3.1 — 310 3.1 2.5–3.5MgO 3.6 2850 210 3.6 —Mullite 3.2 1850 140 5.3 3.0–4.0Borosilicate glass 2.3 — 60–70 3.5 0.5–2Soda-lime glass 2.5 — 60–70 8.9 0.5–1


Glass-matrix materials can be considered as a non-crystalline solid withthe frozen-in structure of a liquid. Characteristics of some important varietiesof glass are given in Table 3.7. Glass–matrix materials are polycrystallinematerials having fine ceramic crystallites in a glass matrix. Important glass–ceramic matrix materials are as follows.

∑ Li2O–SiO2 (LAS). The trade names of such glass–ceramic matrix materialsare Corningware, Zerodur and Ceran. This type of glass–ceramic matrixmaterial has nearly zero thermal expansion and high thermal shockresistance. It is used for the production of optical and telescopic mirrors.

∑ MgO.Al2O3.5SiO2. This system is utilized for the production of variousstable crystalline phases. These are (i) 2MgO.2Al2O3.5SiO2 (known ascristobalite, tridymite and cordierite), (ii) enstatite (MgO.SiO2), and (iii)mullite (3SiO2.2Al2O3). These are used for making radar antennae andradomes for aircraft.

∑ SiO2–Al2O3–MgO–K2O–F. The presence of the mica phase in this systemhelps easy machinability of the product.

∑ SiO2–Al2O3–CaO. In this system wollastonite (CaO–SiO2) or anorthite(CaO–Al2O3–2SiO2) is present as main crystalline phase.

3.4.1 Properties of fibre-reinforced glass/glass–ceramicmatrix composites

During the past few years, the interest in composite materials that couldextend this temperature capability has been stimulated by the results ofseveral research programmes dealing with glass-matrix composites reinforcedwith either graphite [67–70] or alumina fibres [71]. In all of these programmesit was found that the use of fibres exhibiting high strength and stiffness wassuccessful in reinforcing lower-modulus glass matrices. The graphite fibre-reinforced glass system demonstrated exceptionally high levels of strength,fatigue resistance and fracture toughness but was susceptible to fibre oxidationduring elevated temperature oxidation. In contrast, the alumina fibre-reinforcedsilica matrix composite [71] was unaffected by exposure to temperaturesabove 1000oC in air; however, the overall levels of strength and toughness

Table 3.7 Characteristics of some important varieties of glass56

Glass Softening point Density Toughness(∞C) r (g.cm–3) KIC (MPa.m1/2)

Soda-lime glass 700 2.4 0.7Borosilicate glass 825 2.3 0.896% Silica glass 1500 2.5 —Fused quartz 1580 2.6 —


attained were far lower than those of the graphite fibre-reinforced glasssystem. The silicon carbide fibre-reinforced glass-matrix composites possessa unique combination of both high levels of mechanical performance alongwith excellent oxidation resistance. The AVCO System Division (Lowell,MA) is now producing a SiC monofilament of 140 mm diameter that isfabricated by CVD onto a carbon filament core in a pilot plant. This fibreexhibits an average tensile strength of up to 3450 MPa, has a temperaturecapability of over 1300∞C, and is stable in oxidizing environments. It has adensity of 3.2 g cm–3 and an elastic modulus of 415 GPa. This fibre is nowavailable on the commercial market. The second type of SiC fibre, recentlysynthesized by Yajima et al. [72] in Japan, consists of continuous length SiCyarn that is produced from an organometallic polymer. The tows of yarncontain 2000 fibres per tow with an average fibre diameter of 10 mm. TheSiC fibre is highly flexible with an extremely smooth surface. The manufacturer,with a use temperature of up to 1300∞C, has reported tensile strengths of upto 3450 MPa for this fibre. The SiC yarn density is approximately 2.7 g cm-3

and the elastic modulus is 221 GPa.

Composite fabrication

The steps in the fabrication of a SiC monofilament-reinforced glass compositeare as follows.

The SiC monofilament is wound with the desired spacing on a drum,bonded together at periodic intervals with polystyrene, cut into individualtapes and then stacked up in the hot-pressing die to form the composite byalternating SiC tapes with 7740 glass powder (7740 is the Corning GlassWorks designation for a borosilicate glass). The amount of powder is variedto make two types of composite samples, one with 65 vol% SiC and theother with 35 vol% SiC fibre. The composite lay-up is then hot-pressed for20 min at a maximum temperature of 1150∞C and a pressure of 6.9 MPa inan argon atmosphere. It is found that this hot-pressing schedule results incomplete densification of the 7740 glass with very little bubble or voidformation. A cross-section of a typical 65 vol% SiC composite is shown inFig. 3.18.

The SiC yarn-reinforced 7740 glass specimens are fabricated using theidentical procedure developed for the fabrication of graphite fibre-reinforcedglass-matrix composites [69, 70]. The SiC yarn is passed through a slurry ofglass powder and isopropyl alcohol, dried and then cut into tapes of theappropriate length to fit the hot-pressing die. Sufficient tape is then stackedin the die to obtain the desired thickness composite and hot-pressed in vacuumfor 1 hour at 1200oC and at a pressure of 14 MPa. This hot-pressing scheduleis found to result in void-free composites with complete densification of theglass matrix. The coefficients of thermal expansion of 7740 glass-matrix


composites are given in Table 3.8. These values were computed using thechange in dimension measured between 22 and 500oC and indicate a relativelyminor anisotropy in thermal expansion as compared with other compositesystems such as graphite-reinforced glass [70]. The properties of SiC fibre-reinforced 7740 glass matrix composites are given in Table 3.9. The mechanicalproperties of freeze-gelled, unidirectional carbon-fibre-reinforced CMC aregiven in Table 3.10 and the mechanical properties of some fibre-reinforcedsol-gel glass matrix composites are given in Table 3.11. From the fracturecharacteristics of the silicon carbide yarn glass-matrix system, it appears thatthe crack blunting ability of this system at its present state of developmentis somewhat less than that for the silicon carbide monofilament glass system(Fig. 3.19).

3.18 Cross-section of 65 vol% SiC monofilament-reinforced 7740borosilicate glass composite (reproduced by permission of MatrixComposite Mat. Sci.)66.

Table 3.8 Code 7740 glass-matrix composite coefficients ofthermal expansion, CTE (average value between 22∞C and500∞C)66

Filament Orientation CTE(10–6 oC)

35 vol% SiC monofilament 0o 4.2090o 4.60

40 vol% SiC yarn 0o 3.2590o 2.70

50mm


Table 3.9 Properties of SiC fibre-reinforced 7740 glass-matrix composites66

Monofilament Yarn

Fibre content (vol%) 35 65 40

Density (g.cm–3) 2.6 2.9 2.4

Axial flexural strength (MPa)at 22∞C 650 830 290at 350∞C — 930 360at 600∞C 825 1240 520

22oC axial elastic modulus (GPa) 185 290 120

Axial fracture toughness (MN.m–3/2)at 22∞C 18.8 — 11.5at 600∞C 14.3 — 7.0

Table 3.10 Mechanical properties of freeze-gelled, unidirectional carbon-fibre-reinforced CMC made by hand lay-up73

Sample Density (¥103) Flexural Strain at Dynamic(kg m–3) strength peak load modulus

(MPa) (%) (GPa)

Without added 1.6 (0.02)† 118 (17) 0.4 (0.05) 40 (4)CDM 105*glass-ceramic

With added 1.7 (0.01) 212 (10) 0.7 (0.15) 43 (5)CDM 105*glass–ceramic

*CDM 105 is a proprietary glass–ceramic, manufactured by the freeze gelation method.†Figures in parentheses are standard deviations.

Table 3.11 Mechanical properties of some fibre-reinforced sol-gel glass-matrixcomposites73

Fabrication Fibre Final density Dynamic Flexural Work ofroute (¥103 kg m–3) modulus strength fracture

(GPa) (MPa) (KJ m–2)

Cast Saffil 2.06 53.1 25 0.16Hand lay-up Carbon Mat 1.72 42.6 212 —Filament Nextel 2.26 — 179 3.10Wound F P Alumina 2.19 — 210 1.70

Carbon 1.60 36.0 270 13.30


3.4.2 Glass-ceramic matrix composites

A range of glass-ceramics, each having different crystalline phases, has beenused as the matrix material. The most widely used matrices are based on thelithium aluminosilicate (LAS) system. The constituents of typical glass–ceramic matrices are given in Table 3.12.

The strain to failure of SiC fibres is greater than that of the glass–ceramicmatrix. During loading on silicon carbide fibre-reinforced glass–ceramiccomposite, matrix cracking occurs before failure of the fibres. The load isbeing transferred from the matrix to the fibres. The room-temperature modulusvalue of LAS is less than 100 GPa compared to over 200 GPa for the SiCfibres. It is therefore considered that there is potential for increasing thestiffness of the glass–ceramic matrix by incorporating SiC fibres. The tensilemodulus data of LAS–SiC composites made with both unidirectional andcross-plied SiC fibres are presented in Table 3.13.

Compared with other structural materials there is a relative dearth offatigue data for ceramics, and it is therefore not surprising that little isknown of the fatigue behaviour of glass–ceramic matrix composites. It appears

20 mm

3.19 Fracture surface of SiC yarn-reinforced 7740 borosilicate glass(reproduced by permission of Matrix Composite Mat. Sci.)66.

Table 3.12 Typical compositions of lithium aluminosilicate (LAS) glass–ceramics74

LAS system Constituents

LAS I Li2O–Al2O3–MgO–SiO2 with addition of ZnO–ZrO2 and BaO

LAS II Li2O–Al2O3–MgO–SiO2 with addition of Nb2O5, ZnO, ZrO2 and BaO

LAS III Li2O–Al2O3–MgO–SiO2 with addition of Nb3O5 and ZrO2


that at room temperature the fatigue performance is quite good; in tension–tension fatigue the maximum fatigue stress has to exceed about 0.7 of thetensile fracture stress to cause failure within 105 cycles in any of theunidirectional LAS matrix composites (Table 3.14). Residual strengthmeasurements on specimens which survive 105 cycles suggest that significantdamage, and hence a reduction in residual strength, occurs mainly when themaximum fatigue stress is greater than the stress required for matrix micro-

Table 3.13 Modulus data for LAS–SiC composites74

Matrix Vol. % Tensile modulussystem SiC (GPa)

LAS 0 46LAS I 46 (unidirectional) 133LAS II 46 (unidirectional) 130LAS II 44 (unidirectional) 136LAS I ~50 (cross-plied) 118LAS III ~50 (3-D braid) 79–111

Table 3.14 Fatigue behaviour of SiC fibre-reinforced LAS74

Material Tensile Max. fatigue Fatigue Residualstrength, s stress, smax

ssmax cycles strength

(MPa) (MPa) (MPa)

LAS I 261 207 0.79 1.9 ¥ 102 —(unidirectional) 261 207 0.79 2.2 ¥ 103 —

261 172–138 0.66–0.53 >105 286–226

LAS II 550 355 0.65 >105 485(unidirectional) 550 310 0.56 >105 525

550 275 0.50 >105 485550 225 0.41 >105 620

LAS III 575 456 0.79 5 ¥ 101 —(unidirectional) 575 421 0.73 5 ¥ 102 —

575 357 0.62 >105 462575 315 0.55 >105 480575 280 0.49 >105 602575 223 0.39 >105 538

LAS II 325 361 1.10 6 ¥ 101 —(cross-plied) 325 285 0.88 1.2 ¥ 102 —

325 280 0.86 1.5 ¥ 102 —325 280 0.86 3 ¥ 102 —325 262 0.81 3 ¥ 102 —325 256 0.79 3 ¥ 104 —325 245–210 0.75–0.65 >105 385–315

LAS III 269 181 0.67 104 —(cross-plied) 269 179 0.67 3 ¥ 103 —

269 179 0.67 5 ¥ 103 —269 186–163 0.69–0.61 >105 291–227


cracking. This is evident from the data of Table 3.14 bearing in mind thatLAS-1 exhibits a linear stress–strain curve to failure, i.e. no micro-cracking,whereas the matrix cracking stress for LAS II and LAS III is about 270 MPaand 320 MPa respectively [74].

3.4.3 Silicon carbide whisker-reinforced glass andglass–ceramic composites

Several-fold improvements in strength and fracture toughness of glass andglass–ceramics have been obtained by reinforcing these matrices with siliconcarbide whiskers. Incorporation of whiskers in the matrices results in atremendous increase in viscosity, thus necessitating much higher compositeprocessing temperatures compared to that of the matrix alone. Glass andglass–ceramic composites eliminate the residual glassy phase problem (whichis creating the thermal expansion mismatch) and the increasing refractorinessof the matrix, thus enhancing the high-temperature properties. Several glass–ceramic compositions were studied [10] as matrices for whisker reinforcement:see Table 3.15.

The first composite system is a barium osumilite composition reinforcedwith 30 wt% (25 vol%) No. 1 whiskers. The second composite system evaluatedis a barium-stuffed cordierite matrix 30 wt% No. 1 whisker composite. Theroom-temperature properties of these two composites are given in Table3.16.

Figure 3.20 shows the change in modulus of rupture (MOR) and fracturetoughness of a Ba–osumilite composite system with temperature [10]. Theroom-temperature properties are maintained to 900oC. Beyond 900oC theproperties begin to deteriorate. The flexural strength drops from 406 MPa atroom temperature to 55.1 MPa at 1200oC. There is an increase in KIC beyond900oC. The fracture toughness value peaks at 1000oC and decreases again.The small amount of residual glass affects the composite performancesignificantly beyond 900oC. Figure 3.21 shows the change in MOR andfracture toughness of a Ba-stuffed cordierite composite system withtemperature. The room-temperature properties are maintained to 900oC. Beyond

Table 3.15 Glass–ceramic composition range10

Ba–stuffed cordierite Ba–osumilite

4MgO.4Al2O3.10 SiO2 2BaO.4 MgO.6Al2O3.18SiO2

BaO2+ + 2Al3+ = 2Si4+ x [Ba2+ + 2Al3+] = 2x Si4+

x BaO.4 MgO.(4 + x)Al2O3 n [Mg2+ + Si4+] = Al3+

¥ (10 – 2x) SiO2

Where x = 0 to 0.5 (2 – x)BaO.(4 + n) MgO.(6 – x – n) Al2O3

¥ (18 + 2x + n) SiO2

where x = 0 to 0.5, n = 0 to 1.0


Table 3.16 Properties of Ba–osumilite and Ba-stuffed cordieritematrix composites containing 30 wt% No. 1 whiskers10

Property Ba–osumilite Ba–stuffedmatrix cordierite

matrix

MOR (MPa) 400 358KIC (MPa.M1/2) 4.5 4.5Young’s modulus (GPa) 156.4 186Shear modulus (GPa) 62.0 73Poisson’s ratio 0.262 0.274Thermal expansion at 35.5 36.225–1000∞C (¥10–7/∞C 2808 2770Density (kg/m3)

25 900 1000 1100 1200Temperature (∞C)

MOR

KIC

420

350

280

210

140

70

MO

R (

MP

a)

7

6

5

4

3

2

1

KIC

(M

Pa.

m

)

3.20 Variation of properties of Ba–osumilite matrix composite withtemperature (reproduced by permission of Am. Ceram. Soc.Bull.)10.

25 900 1000 1100 1200Temperature (∞C)

MOR

KIC

350

280

210

140

70

MO

R (

MP

a) ●

5

4

3

2

1

KIC

(M

Pa.

m

)

3.21 Variation of properties of Ba-stuffed cordierite matrix compositewith temperature (reproduced by permission of Am. Ceram. Soc.Bull.)10.


900oC the properties begin to deteriorate. The composite strength decreasesfrom 358 MPa to 213.6 MPa at 1200oC. A much larger fraction of room-temperature strength (60%) is thus retained to 1200oC compared to the Ba–osumilite system (13%). Figure 3.22 shows SEM micrographs of the fracturesurfaces of Ba–osumilite composite tested at room temperature and at 1100oC.The flow of the glass is quite evident on the 1100oC fracture surface.

SiC whiskers are available from various sources. Table 3.17 gives thecomposition of the glass matrices. Code 0080 (a soda lime glass), code 1723(an aluminosilicate glass), code 7052 (a borosilicate glass) and code 7740(another borosilicate glass) were used as matrices. Figure 3.23 shows theeffect of the volume fraction of SiC whiskers in composite with code 1723matrix on the critical aspect ratio, according to Dow’s analysis and Rosen’sanalysis. Figure 3.24 shows the effect of porosity on flexural strength ofcomposites with code 1723 glass matrix [75, 76].

Figure 3.25 shows a SEM micrograph of a well-consolidated 30 wt% SiCwhisker-reinforced composite [10]. Table 3.18 gives the mechanical properties

(a) (b)

3.22 Fracture surface of Ba–osumilite matrix composite tested at (a)25oC and (b) 1100oC (bar = 1 mm) (reproduced by permission of Am.Ceram. Soc. Bull.)10.

Table 3.17 Composition of glass matrices10

Oxide Amount present (wt%)

Code Code Code Code0080 1723 7052 7740

SiO2 73 57 64 81Al2O3 1 16 8 2B2O3 — 4 19 13Li2O — — 1 —Na2O 17 — 2 4K2O — — 3 —MgO 4 7 — —CaO 5 10 — —BaO — 6 3 —


0 10 20 30 40Volume fraction

140

120

100

80

60

40

20

0

(l/d

) C

3.23 Effect of volume fraction of SiC whiskers in composite with code1723 matrix on critical aspect ratio, according to (�) Dow’s analysisand(�) Rosen’s analysis (reproduced by permission of Am. Ceram.Soc. Bull.)10.

10 30 40% porosity

350

280

210

140

70

0

Flex

ura

l st

ren

gth

(M

Pa)

50 wt%30 wt%

40 wt%

3.24 Effect of porosity on flexural strength of composites with code1723 matrix (reproduced by permission of Am. Ceram. Soc. Bull.)10.

3.25 Fracture surface of composite containing 30 wt% No. 1 whiskersin code 1723 matrix (bar = 10 mm) (reproduced by permission of Am.Ceram. Soc. Bull.)10.


and aspect ratio data on all these whiskers as available from the manufacturers[79].

SiC whisker-reinforced glass matrix composites were fabricated at thesame process viscosity of the matrices and were well consolidated. All thecomposites were 30 wt% whiskers (No. 1) composites. The properties ofthese composites are given in Table 3.19. A comparison of the 1723 matrixcomposite and the 7052 composite shows that the latter is much weaker andhas a lower modulus. Comparing the 7052 and 7740 systems, the 7740composites are weaker still. A comparison of the 0080 and 1723 systemsagain shows a lower performance for the 0080 composite.

The experiments were done [10] with the 1723 glass matrix and No. 1,No. 3 and No. 4 whiskers. Table 3.20 summarizes the results of these

Table 3.18 Properties of silicon carbide whiskers10

No. Type of whisker Dimensions (mm) Mechanical properties

Diameter Length Strength Modulus(MPa) (GPa)

1 a + b mixture 0.6 10–80 689 6.892 a + b mixture 0.6 10–80 689 6.893 a + b mixture 1–10 20–400 551 184 a + b mixture 0.05–0.2 10–40 482 205 b–SiC Tangled woolly whiskers No data No data

Table 3.19 Properties of SiC whiskers-reinforced glass matrix composites10

Glass matrix Matrix MOR Composite Matrixexpansion (MPa) modulus (GPa) modulus (GPa)(10–7/∞C)

Code 1723 52 337.6 (± 8%) 141.6 (± 7%) 86 (± 5%)Code 7052 53 241.1 (± 8%) 107.5 (± 6%) 57.8 (± 8%)Code 7740 34 194.3 (± 10%) 92.3 (± 8%) 62.6 (± 8%)Code 0080 95 179.1 (± 9%) 96.4 (±10%) 70.3 (± 7%)

Table 3.20 Comparison of whisker performance in code 1723 glass matrixcomposites10

Composite Whisker Whisker MOR Modulus Fractureno. no. content (MPa) (GPa) toughness

of composite(wt.%) (MPa. m1/2)

1 1 30 338 (± 8%) 141.6 (± 5%) 3.4 (± 7%)2 4 30 193 (± 10%) — —3 3 30 263 (± 8%) 133.7 (± 6%) 2.1 (± 10%)4 4 30 190 (± 11%) — —5 4 30 250 (± 8%) 128.8 (± 5%) 2.8 (± 8%)


experiments. All the experiments were done at 30 wt% (25 vol%) whiskerloading, except for composite 4, which contained 20 wt% whiskers. Thecomposite consolidation was conducted under the same conditions forcomposites 1 through 4, while composite 5 was processed at a highertemperature. The flexural strength measurements show that the No. 1 whiskersgive the best strength values. The composite containing No. 1 whiskers ismuch stronger (338 MPa) than the No. 3 composite (263 MPa). The modulusof the No. 3 whisker composite is also lower (133.7 GPa) compared to theNo. 1 composite (141.6 GPa).

3.5 Microstructural observation

Composite samples are sectioned with a diamond saw and mounted in coldcuring epoxy resin. Because of their porous nature, the composites are infiltratedunder vacuum and subsequently cured under pressure in order to force themounting resin into the pores. Mounted samples are ground flat on 240grit silicon carbide paper, finely ground with a 9 mm oil-based diamondslurry and finally polished with a 1 mm diamond slurry and a 50 nm silicasuspension.

3.5.1 Scanning electron microscopy (SEM)

The polished samples are sputtered with a thin layer of gold for analysis ina scanning electron microscope (SEM), a Jeol JSM 35c fitted with a link AN10000 energy-dispersive X-ray spectrometer (EDS). The fractured surfacesand polished sections through fractured specimens can also be prepared andanalysed in this manner. SEM analysis may reveal a non-uniform fibredistribution in the composite. In composites sintered at different temperatures,cracking in the matrix phase and residual porosity can be identified and thefiller particles are discernible. The EDS indicates the higher particles and thematrix constituents.

3.5.2 Transmission electron microscopy (TEM)

TEM analysis is performed in a Jeol 2000 FX equipped with an EDS system.Thin sections (approximately a few hundred nanometres thick) suitable forTEM are prepared by cutting 3 mm slices, grinding them to a thickness of~300 mm and dimpling them to leave a central region ~10 mm thick. Thinningwith argon ion bombardment in a Gatan Duomill may be carried out untilspecimen perforation occurs. TEM studies of the composite sintered at differenttemperatures may reveal the possibility of forming a fused network with thespherical particles originally present in the samples. The filler particles may


exhibit strain bands and the matrix may contain several twinned regions,both suggestive of residual strains within the material.

3.5.3 X-ray diffraction (XRD) analysis

XRD spectra of composites sintered at different temperatures are obtainedby using a defractometer. Samples sintered at a particular temperature indicatethe structural behaviour as either amorphous or crystalline in nature. Suchindication is of immense help to researchers for improving the properties ofthe resultant materials by optimizing the rate of sintering temperature.

3.5.4 Optical microscopy analysis

An extensive network of porosity of the sintered samples is obtained byusing an optical microscope. Dark-field illumination in the optical microscopereveals the pores to have a crystalline surface texture and readily distinguishesbetween the denser (darker) and more porous (lighter) regions of the specimens.Also, a degree of subsurface detail is revealed. Optical microscopy is carriedout in reflected light with Nomarski differential interference contrast (DIC)and dark-field modes on a suitable microscope. The fibre volume fractioncan be estimated for all the samples. Matrix cracks around fibres arisingfrom residual stresses can be observed in all samples and tend to reach thespecimen surface via dense matrix regions. Circumferential crack patternsmay indicate residual stresses arising from fibre/matrix thermal expansionmismatches.

3.6 Application areas

Very important application areas of glass/glass–ceramic matrix compositesare supersonic planes, US high-speed civil transport planes, turbine blades,heat shields, rocket nozzles and propulsion components, re-entry thermalprotection for spacecraft, rocket cone frustra, braking materials, cutting toolinserts, high-wear parts such as wire drawing or extruding dies, valve seats,high-precision balls, bearings for corrosive environments, and plungers forchemical pumps. Other application areas for glass/glass–ceramic matrixcomposites are disc brakes for racing cars and aircraft, gas turbine components(e.g., exhaust nozzle flaps and seals), nose cones and leading edges formissiles, and biomedical implants such as bone plates.

The applications of glass/glass–ceramic matrix composites (CMC) can bedivided into two specific categories: aerospace applications and non-aerospaceapplications. In aerospace applications, performance is the prime consideration,while in non-aerospace applications cost-effectiveness is paramount. Thecharacteristic properties of materials for aerospace applications should be


high strength-to-density ratio, high stiffness-to-density ratio, improved damagetolerance at significantly higher temperatures and/or faster cruising speeds,and improved flight performance at higher altitudes. Continuous fibre-reinforced ceramic matrix composites potentially offer higher specific strengthproperties, which can be utilized in various high-temperature aerospaceapplications. Silicon carbide coated carbon/carbon composites and carbon/silicon carbide composites are the right candidate materials for such high-temperature aerospace applications [65, 66]. The applicable temperature rangeof CMCs is 800–1650∞C. The tensile and compressive strengths of CMCsare in the range of 175–350 MPa and their moduli in the range of 100–175GPa throughout the temperature range of 800–1650oC.

In non-aerospace applications of CMCs, cutting tool inserts, wear-resistantparts, energy-related applications such as heat exchanger tubes, nozzles,exhaust ducts, etc., are the emerging areas. Particle and whisker-reinforcedceramics are commonly used for cutting tools, wear-resistant parts and heatengine applications. TiC particle-reinforced Si3N4 and Al2O3 and SiC(w)–Al2O3 composites are used for cutting tool inserts. Toughened zirconia andSiC whisker or continuous fibre-reinforced composites are used for makingwear-resistant parts. The other non-aerospace application areas of carbon–carbon composites are biomedical implants and internal fixation of bonefractures because of their excellent biocompatibility. They are also used formaking moulds for hot-pressing. Carbon–carbon moulds can withstand higherpressures and offer a longer life than polycrystalline graphite and tool steelmaterial. In general, their high cost limits applications to aerospace andother special applications.

Whisker-reinforced glass–ceramic matrices are expected to find severalapplications in automotive components, metal forming, cutting tools, etc., dueto their low thermal expansion, high thermal shock resistance, high reliabilityand low material and processing costs. Some industrial applications for continuousfibre-reinforced ceramic matrix composites (CMCs) are listed below.

∑ Heat engine liners, combustors, high-wear parts, etc. CMCs are used inhigh-temperature gas turbines.

∑ CMCs are used in the manufacture of preheaters and recuperators in heatrecovery equipment. They are used for indirect heating and energy-intensiveindustrial internal processes such as glass melters, steel reheaters andaluminium remelters.

∑ Radiant tube burners are made by using CMC. These are used for indirect-fired, high-temperature zones, controlled atmosphere heating and meltingapplications.

∑ Reformers and reactors of chemical process equipment.∑ Handling equipment, internals and cleanup of waste incineration systems.∑ Filters, substrates and centrifuges of separation and filtration systems.

Such filtration and separation systems can be used for gas turbines,


particulate traps for diesel exhausts, molten metal filters and sewagetreatment equipment.

3.7 Future trends

The gel synthesis of ceramic powders offers many advantages over moreconventional methods based on solid-state reactions: lower processingtemperature, better control of the morphology and microstructure, powderlessprocessing of ceramics. However, a better knowledge of sol-gel chemistryhas to be developed before a real mastery of the process can be reached. Thisrequires the careful characterization of all the chemical species formed duringthe course of the sol-gel process. The chemical modification of alkoxideprecursors opens future possibilities for the molecular design of advancedceramic matrix composite (CMC) materials. There is a possibility of developingpotential high-strength viable CMC-based engineering materials by increasingtoughness, decreasing sensitivity to flaws and increasing the reliability ofthe materials. The incorporation of carbon fibre into glass and glass–ceramicmatrices may produce composite materials having toughness several timesgreater than that of monolithic matrix materials, which can withstand largetensile strains prior to final failure and can exhibit a significant degree ofdamage tolerance material. Conventional polymer-matrix composite processingtechniques such as sheet moulding compound (SMC), filament winding andresin transfer moulding (RTM) may be utilized for the fabrication of advancedceramic matrix composite prepregs before putting them into the firing schedule.Improvement of the freeze-gelation method may eradicate many problemsassociated with the CMC fabrication process and may develop high-fracture–toughness materials by reducing porosity and internal shrinkage.

3.8 References

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7. Brennan, J.J., Chyung, K. and Taylor, M.P., ‘Reaction inhibited-silicon carbide fibre


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67. Phillips, D.C., Sambell, R.A.J. and Bowen, D.H., ‘The mechanical properties ofCarbon fibre reinforced Pyrex glass’, J. Mat. Sci., 7, 1454 (1972).

68. Levitt, S.R., ibid., 8, 793 (1973).69. Prewo, K.M. and Bacon, J.F., Proc. 2nd Int. Conf. on Composite Materials, Toronto,

Canada (AIME, New York, p. 64 (1978).70. Prewo, K.M., Bacon, J.F. and Dicus, D.L., SAMPE Q., 42 (1979).71. Bacon, J.F. and Prewo, K.M., Proc. 2nd Int. Conf. on Composite Materials, Toronto,

Canada (AIME, New York), p. 753 (1978).72. Yajima, S., Okamura, K., Hayashi, J. and Omori, M., ‘Glass-ceramic matrix

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Woodhead Publishing, Cambridge, UK (1999).75. Dow, N.F., GEC Missile and Space Division, Report No. R635D61.76. Rosen, B., Fibre Composite Materials, ASM, pp. 37–75 (1964). Process – I

99

4.1 Introduction

Ceramic materials have been produced by man for at least 9000 years. Themicrostructures of most traditional ceramics resemble particulate ceramiccomposites in that at least one of the phases present consists of approximatelyequiaxed, discontinuously distributed particles. Although particulate phasesmay be present naturally in the clay used for shaping, for much of the historyof ceramic technology particulates have also been added deliberately as a‘temper’ of quartz, limestone, sand, shell, ‘grog’ (recycled pulverised pottery)or other easily available substances. The function of these particulates intraditional ceramics is usually to give high-temperature strength so that theshape is retained during firing or to act as a cheap filler, and thus has littlerelevance to this publication. There is evidence, however, that variations inthe choice of temper occurring over periods of many years in particularcommunities resulted in improvements in mechanical properties such asstrength, toughness or thermal shock resistance [1] and point to at least anaccidental application more than 1000 years ago of some of the principlesdescribed in this chapter.

The development of ‘engineering ceramics’ with sufficient strength forapplication in load-bearing situations did not begin in earnest until the 1960s.Research concentrated initially on simplifying and refining the coarse andflaw-ridden microstructures found in traditional ceramics. This was verysuccessful, and the outcome was the production of almost pore-free, fine-grained, single-phase ceramics with strengths of several hundred MPa(compared with several tens of MPa for traditional ceramics). Investigationsthen switched to methods of increasing the toughness. The range of innateceramic toughness values is relatively limited, although the use of ceramicssuch as carbides and nitrides with strong covalent bonds and therefore inherentlyhigh surface energy succeeded in pushing strengths towards 1 GPa. Moresubstantial increases in fracture toughness have proved possible only bymaking the microstructure more complex again, though this time in a systematic

4Particulate composites

R I T O D D, University of Oxford, UK


and targeted fashion, particularly by making many of the types of compositeand related structure described in this book.

The particulate composites described in this chapter constitute perhapsthe simplest departure from a fine-grained single-phase ceramic. Theparticulates do not provide the highest strengths or the greatest degree oftoughening to be found in ceramic composites, but against this they arerelatively cheap and easy to process compared with other shapes ofreinforcement. Particulate reinforcements also provide inherently isotropicproperties (cf. long-fibre composites) and are less toxic and easier to handlethan whiskers. Although particulate composites already see commercialapplication, their simplicity makes them useful as model materials as well,easier to understand and simulate than more complicated microstructures.

This chapter aims to describe the principles behind the processing,microstructural development and properties of particulate ceramic compositesand to illustrate these using experimental results. The main emphasis is onexamples where the addition of particulates to a ceramic matrix causes newmechanisms to operate that give an improvement in properties greater thanwould be expected from a ‘rule of mixtures’. The chapter concentrates almostexclusively on structural composites, since this is where most work has beendone to date. Particulate ‘nanocomposites’ are included in the chapter, sincethe important examples described are currently at the coarse end of the‘nanoscale’, and the principles underpinning their properties seem to be asimple extension of those relevant to the ‘microcomposites’ with which therest of the chapter is concerned.

The next section describes the processing and microstructural developmentof particulate composites, and is followed by a section on thermal residualstresses. These stresses are often the most obvious consequence of addingsecond-phase particles to a matrix and can have a profound effect on properties.Factors determining the toughness, strength and wear resistance of particulatecomposites are then considered in turn, and the chapter concludes with anassessment of possible future developments in this area.

4.2 Powder processing and microstructural

development

Unlike fibre- or whisker-reinforced composites, particulate composites havethe advantage of being compatible with conventional powder processing,and in many cases can be pressurelessly sintered. As with other ceramicmicrostructures, a myriad of other ingenious fabrication routes have alsobeen reported, but these are too numerous and system-specific to describehere. This section merely outlines the main points of powder processingwhere the production of composites in chemically compatible systems (i.e.those in which the components do not react chemically with one another)differs from that of monolithic ceramics.

Particulate composites 101

We begin with milling and dispersion of the powders in a liquid. Inaddition to the role of breaking down hard agglomerates, as for monolithicceramics, this step must also thoroughly mix the component powders of thecomposite. For composites in which the particulates need to be relativelylarge, however, it is important not to reduce the mean size of the particulatesby using a milling treatment that is too aggressive or very long in duration.More careful control of the milling procedure is often required than formonolithic ceramics.

The particulate and matrix powders usually have differing surface responsesin the development of electrical double layers or adorption of steric dispersantsduring milling. This does not usually degrade their ability to mix. It mighteven help it if, for example, the components develop opposite surface charges,though this might affect the viscosity and achievable solid loading of theslurry. Problems arise mainly when one of the components shows no responseto these deagglomeration mechanisms. In these cases the remedy is the sameas for monolithic processing, namely to try different liquids or dispersants,or to coat the particles with a substance that does offer a response.

The other principal differences between monolithic ceramics and powdercomposites occur during sintering. When a particulate second phase that isconsiderably larger than the matrix powder is incorporated into the greenbody, it represents a region that will not shrink with the matrix as sinteringtakes place. The resulting mismatch in shrinkage inhibits sintering of thematrix and can also lead to stresses sufficient to cause cracking [2]. Thediffusional fluxes during sintering can also relax the stresses in the matrix[3], however, essentially through simultaneous diffusion creep. This can besufficient to enable sintering to proceed to completion and for cracking to beavoided. There are many examples of pressurelessly sintered compositescontaining relatively large particles [4–7].

When the particulate phase is smaller than, or of comparable size to thematrix powder, this source of inhibition does not arise. If the particulatephase has similar diffusional properties to the matrix at the sinteringtemperature, sintering can actually be improved because the particles opposegrain growth by pinning the grain boundaries [8]. Examples of this type ofcomposite include Al2O3–ZrO2 and Al2O3–Cr3C2. The ability of the particlesto participate in diffusion usually means that they are mobile. Grain growthis therefore not entirely prevented and the particles are dragged around bythe migrating grain boundaries, coalescing in the process, so that typicalfinal microstructures are characterised by rounded particles of equilibriumshape, often coarser than the original particles added as powder, andpredominantly situated at grain boundaries and (especially) triple lines.

The addition of fine particles that are much more refractory than thematrix has a different effect. The outstanding example of this behaviour is inalumina/SiC ‘nanocomposites’ [9]. The SiC/Al2O3 interface may have a low


diffusion coefficient, and it is widely accepted [10, 11] that it may be difficultto remove matrix material from the particle matrix interfaces, so that theessential sintering process of removal of material from the grain boundariesand deposition in the pores is severely impeded. Figure 4.1 shows that theaddition of only 2.5 wt% SiC severely impedes sintering, and further additionscontinue to reduce the sintered density. Another consequence of the inabilityof the SiC/Al2O3 interface to participate in diffusional processes is that theparticles are immobile. They therefore retain their original size, are uniformlydistributed and are situated both within the grains and on the grain boundaries.They pin grain growth more strongly than mobile particles, approximately inaccord with Zener’s theory [12, 13].

This striking property of the Al2O3/SiC interface can be understood interms of the observation of Ashby and Centamore [14] that the more refractoryof two phases at an interface (the covalently bonded SiC in this case) controlsthe interface reaction because in general atoms in both phases must be involvedin the reaction. The majority of the Al2O3/SiC interfaces in the nanocompositeshave been observed to be free of any glassy phase, the presence of whichwould presumably allow alumina to be removed or deposited at the interfacewithout the involvement of the SiC, and consequently much more rapidly.The introduction of an interfacial layer may be the source of the ability ofsintering aids such as Y2O3 to enable these materials to be pressurelesslysintered [15, 16] (Fig. 4.2).

In conclusion, particulate composites are more difficult to process usingpowders than monolithic ceramics, but are easier than other kinds of compositenevertheless. They can often be sintered to full density without pressure.When this is not possible, sintering aids or the superimposition of pressure(hot pressing, hipping) can be used to alleviate the problems, and there aremany examples of this in the literature and in commercial practice.

0 2 4 6 8 10wt% SiC

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4.1 Sintered density (1700∞C, 2 hours) against SiC content foralumina/SiC nanocomposites, demonstrating the inhibition ofsintering caused by the SiC particles [10].


4.3 Thermal microstresses

If the matrix and particles in a composite have different thermal expansioncoefficients then thermal microstresses develop during cooling from processingtemperatures. These stresses can be very large in particulate ceramic composites,firstly because the processing temperatures are high so that the temperaturechange on cooling is large, secondly because ceramics are typically verystiff so that a large stress develops for a given thermal expansion mismatch,and thirdly because, unlike metals, most ceramic phases have little scope forplastic relaxation of the stresses during cooling, at least below 1000∞C. Thechange of these stresses during cycling of a MgO–SiC ‘nanocomposite’ fromroom temperature to 1550ºC and back again is demonstrated in Fig. 4.3 [17]from which it can be seen that the stress level in the SiC particles is almost4000 MPa at room temperature.

For most particulate composites the mismatch between the particles andthe matrix is more important than the anisotropy of either component (thoughalumina/aluminium titanate composites provide a notable exception and aredescribed below). The main features of the stresses can therefore be understoodin terms of a simple elastic model assuming thermoelastic isotropy andconsisting of a spherical particle in a concentric spherical shell of matrixwith dimensions chosen to give the appropriate volume fractions. The particlesare predicted to be under a uniform hydrostatic stress, sp, after cooling. Ifthe particles have a larger thermal expansion coefficient than the matrix, thisstress is tensile, and vice versa. For small particle volume fractions the stress

1650∞C1600∞C1550∞C1500∞C1450∞C

0.0001 0.001 0.01 0.1 1 10% Yttria

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4.2 Effect of yttria additions on sintered density of alumina–2% SiCnanocomposites for various sintering temperatures (data from [16]).The ‘0.0001%’ points indicate no added yttria.


state in the matrix immediately outside the particle is purely deviatoric witha radial stress of sp, and hoop stresses of –sp /2. These matrix stresses fallaway rapidly as 1/r3 where is r is the radial distance from the particle centre.For higher volume fractions of particles, a uniform hydrostatic image stressis superimposed on these stresses, of opposite sign to the particle stress.

Thermal microstresses influence the fracture of composites. Davidge andGreen [18] showed the ability of the stresses to deflect cracks using glassmatrices with a range of expansion coefficients containing large thoria spheres.The model stress field described above shows that when the particle thermalexpansion coefficient is larger than that of the matrix, the radial stress closeto the particles is tensile so that cracks are deflected away from the particles.With the reverse expansion mismatch, cracks are attracted towards the particles.This deflection of crack path is of relevance to the apparent toughness of thecomposite and to the formation of microstructural elements bridging thecrack interfaces. Even on a planar crack path, the fluctuations between tensileand compressive thermal stresses influence the propagation of the crack,which is impeded when passing through the compressive regions. A furthereffect of thermal stresses is that when combined with the crack tip stressfield they encourage the formation of a microcracked zone around the crack,and the accompanying dilatation can shield the crack tip from the externallyapplied stress. These toughening mechanisms are described in more detail inthe next section.

As well as producing these broadly beneficial effects, thermal microstressescan also degrade the strength of composites. The tensile components ofstress can help in crack initiation. In a composite with a uniform distributionof particles, the tensile components act only over distances comparable with

Error = 0.52 GPa

Heating

CoolingPredicted

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– 2

– 3

– 4Ave

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e h

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4.3 Hydrostatic stress in the SiC particles in a magnesia–10%SiCnanocomposite during a thermal cycle to 1550∞C. Note the smallamount of relaxation during the cycle and the good agreement withthe prediction of a simple elastic model for the stresses [17].


the particle spacing, but non-uniform distributions of particles can lead tomean tensile stresses over considerably greater distances, comparable withthe scale of the local volume fraction variations. The most obviously deleteriouseffect of the thermal stresses, however, is the possibility that they are sufficientlylarge to cause spontaneous microcracking during cooling from the processingtemperature. Davidge and Green [18] first pointed out that since the storedelastic strain energy scales with the particle volume, and the energy requiredto create new crack surfaces scales with its area, there is a critical particlesize for spontaneous cracking below which there is not enough energy availableto allow a crack to form. This is why the MgO/SiC nanocomposite containingthermal stresses approaching 4 GPa described above did not crackspontaneously, despite the fact that the microstresses were considerably largerthan the macroscopic strength of the material. A detailed study of Al2O3–20%SiC composites [19] showed that there was an abrupt transition fromnegligible microcracking with small SiC particles (Fig. 4.4(a)) to generalcracking of the alumina matrix when a mean particle size of 10 ± 3 mm wasexceeded (Fig. 4.4(b)). Figure 4.4(b) shows that the microcracks run radiallyfrom the particles as is consistent with the tensile hoop stresses in thissystem (aSiC < aalumina). This behaviour was explained by a fracture mechanicsmodel which correctly predicted the critical particle size, and showed thatonce nucleated, microcracks can propagate from one particle to the next sothat microcracking can spread throughout the material from a small numberof nucleation sites when the median particle size exceeds this critical value.

4.4 Toughening

One of the primary motivations for the deliberate addition of second-phaseparticles to a ceramic matrix is to increase its toughness. If the particles aretougher than the matrix then the crack resistance energy, R, will be increased,approximately according to the rule of mixtures if the crack simply passesthrough the particles and the difference in toughness between the particlesand the matrix is relatively small. This is obviously of limited value, sincethe composite cannot exceed the toughness of the particles. The compositeapproach is much more powerful if it causes new mechanisms to operate thateither do not occur or are weak in single phase materials. The followingtoughening mechanisms have been investigated for non-transformingparticulate composites.

4.4.1 Crack deflection

All else being equal, a toughening effect occurs if the crack tilts or twistsaway from a planar geometry because this reduces the net crack drivingforce. In homogeneous materials such as glass, cracks tend to propagate in


a planar fashion for the same reason, but non-uniform features such as weakinterfaces and residual stresses can lead to such a deflection in other materials.These may occur in single-phase polycrystals, but there is scope foraugmentation of the effect in composites, and particles with higher stiffnessthan the matrix can also lead to deflection. According to the purely geometricalanalysis of Faber and Evans [20], spherical particles are less effective thanplate- or rod-shaped particles and lead to a maximum increase in apparenttoughness, Kc, of around 30%. Even this modest increase represents anoverestimate of the effect, however, since it does not take into account the

(a)

(b)

4.4 (a) Backscattered SEM micrograph of unetched alumina–3 mm SiCcomposite. The straight lines on the specimen surface are scratchesfrom metallographic preparation. (b) Unetched alumina–13 mm SiCcomposite showing radial microcrack network in the matrix [19].


reason for the deflection, which occurs because it is easier for the crack todeflect than to propagate in a planar fashion. This may be because the non-planar crack path has a lower toughness (weak interfaces), because the residualstress provides extra driving force if the crack is deflected (thermal expansionmismatch), or because the strain energy release rate is greater in the directionof deflection (stiff particles), the implication being that the measured toughnesswould be higher if the crack remained planar. A complete argument shouldconsider the driving force required to cause crack propagation at every pointon the crack path, but it is clear that although crack deflection is importantin understanding the net toughness of a composite exhibiting this effect, it isnot itself a potent toughening mechanism.

4.4.2 Crack bridging

If intact or interlocking ligaments remain behind the advancing crack front,the restraining force they exert reduces the stress intensity at the crack tip,causing an increase in the macroscopically measured toughness. Because thebridges accumulate behind the crack front, the toughening effect increases asthe crack propagates, a phenomenon known as R-curve behaviour. Crackbridging is a very potent toughening mechanism in long-fibre compositesand operates in a similar manner with whisker reinforcements [21]. Thesereinforcement geometries are particularly conducive to crack bridging, butthe mechanism can also operate in less favourable situations. Crack deflectionalong weak interfaces can lead to bridging through geometrical interlockingand causes toughening in monolithic alumina exhibiting intergranular fracture[22, 23]. The presence of particulate reinforcements can enhance this effect.If a particle is to act as a bridge, the key requirement is that the crack pathmust be deflected around its periphery and in doing so tilt or twist through90º or more to form an interlocking section. The main factors determiningwhether or not this is possible are (i) the relative toughnesses of the matrix,the particle and the interface, (ii) the residual stress state around the particle,and (iii) the size of the particles. It is important that the interface is relativelyweak. If, for example, the interface and matrix are as tough as the particle,the crack will go through the particle instead of around it. If the particle istough but the interface is only marginally weaker than the matrix, the crackwill tend to detach from the particle instead of undergoing the severe deflectionrequired for interlocking to occur [24]. Particles with thermal expansioncoefficients greater than that of the matrix will have tensile stresses acrossthe interface, effectively weakening it, thus favouring bridge formation. Theinfluence of residual stresses on a crack path increases with the distance overwhich the stresses act, so this effect is more important with bigger particles.

The balance between these different considerations can vary considerablybetween different composite systems. Figure 4.5, for instance, shows


indentation crack paths in alumina matrix composites containing 20 vol% of(a) TiN and (b) Cr3C2 particles [25]. The particles are of similar size in thetwo composites, and neutron diffraction measurements of the residual stressesshowed that both types of particle were in tension with a mean stress closeto 400 MPa. Despite these similarities, Fig. 4.5 shows a profound differencein the crack paths in the two composites. The crack tends to cut through theparticles in the alumina/TiN composite (Fig. 4.5(a)), indicating a relativelystrong particle–matrix interface. In contrast, the crack goes around the peripheryof all the particles encountered in the alumina/Cr3C2 composite (Fig. 4.5(b))even when large tilts are required to do so, and this has resulted in extensive

(b)

(a)

4.5 Indentation crack paths in alumina matrix composites containing20 vol% of particles of (a) TiN and (b) Cr3C2. The alumina matrix isthe darker phase in each case.


crack bridging. Fracture of the surrounding matrix has occurred as theinterlocking areas have been pulled apart. This behaviour must be attributedto a particle/matrix interface that is considerably weaker than that in thealumina/TiN composite.

The differing crack behaviour in these materials is reflected in the tougheningincrement relative to pure alumina that the particles produced. The additionof Cr3C2 increased the toughness considerably, DKc ~ 6 MPa m1/2, whilstTiN particles give only a small toughening effect compared with alumina,DKc < 2 MPa m1/2. This demonstrates the ability of crack bridging to causesignificant toughening, even in particulate composites.

4.4.3 Microcrack toughening

Section 4.3 showed how thermal microstresses in particulate ceramiccomposites can cause spontaneous microcracking when the particles exceeda critical size. For composites in which the particles are below the criticalsize for spontaneous fracture, the imposition of additional stress can lead tostress-induced microcracking. A potential consequence of this is thedevelopment of a process zone of microcracked material ahead of the cracktip. The consequent reduction in modulus ahead of the crack tip reduces thestress intensity [26], though this small effect is countered by the reduction intoughness as a result of the microcracking. The energy dissipated in thewake of a propagating crack as the newly microcracked material is unloadedprovides a stronger effect. This originates both in the irreversible dilatationof the material as the microcracks form in a manner analogous to thetransformation toughening of zirconia, and in the accompanying reductionin stiffness.

The existence of a microcracked zone around cracks in SiC–TiB2 composites,in which the thermal expansion mismatch puts the TiB2 particles in a state oftension after sintering, has been confirmed using small-angle X-ray scatteringand by direct observation in the TEM [27]. The toughness was observed toincrease with crack propagation (R-curve behaviour), as would be expectedby this mechanism which, like crack bridging, relies on the development offeatures behind the crack tip. The observations of microcracking were usedto estimate the extent of microcrack toughening expected, and the resultswere of similar magnitude to the measured toughening increments, definedas the difference between the toughness on initial crack propagation and theplateau value at large extensions. Since the observed toughness increaseswere all less than 2 MPa m1/2, however, it is difficult to separate unequivocallythe contribution of microcracking from those of other mechanisms capableof causing R-curve behaviour such as crack bridging, which would also befavoured by the tendency of the residual stress to aid circumferential crackformation.


4.4.4 Direct influence of thermal microstresses

As a loaded crack propagates through a particulate composite, any thermalmicrostresses of the type described in Section 4.3 will contribute to the totalstress intensity at the crack tip. Stresses applied immediately behind thecrack tip make the biggest contribution to the stress intensity, so the crackcan be expected to propagate easily through regions of residual tension, andconversely will be impeded in regions of residual compression. In order forsignificant crack extension to take place, the externally applied stress intensitymust be sufficient to force the crack front through the compressive regions,and the apparent toughness will therefore be higher than it would be withoutthe microstresses. This principle applies regardless of the sign of the thermalexpansion mismatch because the average microstress on a planar surfacethrough the composite must be zero; any region of tension must therefore bebalanced by a region of compression, though the details of the crack geometryat the critical position of maximum resistance can be expected to differaccording to whether the particles are in compression or tension.

A simple estimate for the toughening increment, DKc, attainable by thesemeans can be obtained by assuming that the regions of compression comprisea uniform microstress, s, acting over a distance d behind the crack tip.Ignoring microstresses acting further behind the crack, which is assumed tobe long, gives [28]:

DK dc = 2 2s p (4.1)

The mean compressive stress in the matrix of the composites shown in Fig.4.5 is ~100 MPa, and the distance between particles ~5 mm. Inserting thesevalues into eq. (4.1) gives a toughening increment of only 0.4 MPa m1/2,demonstrating that this mechanism makes only a weak contribution to thetoughness of most particulate composites.

In summary, crack interface bridging provides the most effective tougheningmechanism in non-transforming particulate ceramic composites and canproduce significant toughening increments compared with the unreinforcedmatrix, particularly when combined with several weaker tougheningmechanisms which are also known to operate. The aim of increasing thetoughness is to influence the more directly applicable properties of ceramics,and the next section examines the effect of particulates on their strength.

4.5 Room-temperature strength

4.5.1 Microcomposites

The aim of toughening a matrix by adding second-phase particles is essentiallyto increase its strength. For cracks in homogeneous materials under uniform


loading the toughness, Kc, and strength, sf, are related by the followingexpression:

Kc = Ysf c (4.2)

where c is the critical crack length and Y is a factor depending on the geometryof the crack. Equation (4.2) predicts that the strength should be proportionalto the toughness assuming that the particulate additions do not affect theflaw size. In some cases, this prediction is borne out in practice, as isdemonstrated in Fig. 4.6 for glass–alumina particle composites [29]. Whencombined with the value of Y appropriate for semicircular edge cracks, thegradient of the line in Fig. 4.6 indicates a crack radius of about 100 mm.

In most cases, however, the strength change on the addition of particlesthat increase the measured toughness is less than this naive prediction wouldsuggest, and often the strength is actually reduced. There are several relatedreasons for this. One which is obvious from eq. (4.2) is that the addition ofthe particulate may increase the critical flaw size as well as increasing thetoughness. This may occur because of some of the processing problemsmentioned in Section 4.2, most simply the presence of pores arising from theinhibition of sintering. Non-uniform particle dispersions can also increasethe effective flaw size dramatically. Clusters of particles can fail to sinterproperly if they are refractory and act as critical flaws. Less extreme localvolume fraction variations can also lead to large strength reductions becauseof the possibility of differential shrinkage cracking during sintering. In systemswith a pronounced thermal expansion mismatch between the particles andthe matrix, regions that are particularly rich or deficient in particles lead tothermal stresses that can aid crack initiation.

The preceding reasons for strength reductions are a consequence ofmicrostructural deficiencies, but even ‘perfect’ microstructures, in which thecritical flaw size is the same as for the unreinforced matrix material, canyield smaller strength increases from particle-induced toughening than would

0 0.5 1 1.5 2Toughness (MPa m1/2)

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4.6 Relationship between strength and toughness for glass/aluminacomposites with best fit straight line to eq. (4.2). Data from [29].


be predicted on the basis of eq. (4.2). This is because the most potent tougheningmechanisms, crack interface bridging and microcrack toughening, rely onprocesses that occur behind the crack tip, and consequently exhibit R-curvebehaviour, as described in Section 4.4.2. Taking the example of a particulatecomposite toughened by bridging of the crack by the particles, as an initiallysmall crack grows between particles the toughness is simply that of thematrix. If tensile thermal stresses are present in the matrix, the effectivetoughness at first gets smaller as the crack grows longer because these internalstresses aid crack propagation. After further propagation, however, the cracktip passes a particle or other bridging element which subsequently exerts aclosure force on the crack, causing an increase in the macroscopically measuredtoughness. As crack propagation continues, more bridges are formed and thetoughness continues to rise until the bridges that were formed first are broken,beyond which point a steady state is established and the toughness becomesconstant.

Detailed consideration of the strength under these conditions (see, e.g.,[30]) shows that for small critical flaws, the strength is determined solely bythe toughness versus crack length behaviour in the very early stages of crackgrowth, before any bridges have been formed. Conversely, the strength forvery large flaws is determined essentially by the plateau toughness. Thestrength for intermediate flaws depends on the form of the toughness–cracklength relationship as it rises steeply after formation of the first bridges. Inthis region, some stable crack growth can occur prior to final failure as thetoughness rises more quickly with crack growth than the applied stress intensity,and the strength becomes insensitive to flaw size.

Combining the information in the previous two paragraphs enables thefollowing common observations regarding strength–toughness relationshipsin particulate composites to be explained straightforwardly:

∑ Most standard methods of toughness measurement rely on the use of largeflaws and therefore measure values close to the plateau toughness. Thelack of correlation between measured toughness and strength that isfrequently observed can partly be attributed to the fact that flaws inparticulate composites are usually sufficiently small for the strength to bedetermined by the toughness during initial crack propagation rather thanby the plateau toughness.

∑ If there are no residual stresses in the composite, the strength of samplescontaining only very small flaws can be expected to be similar to that ofthe matrix alone. If tensile stresses are present because of the thermalexpansion mismatch between the matrix and the particles, the initialreduction in apparent toughness during crack growth means that the strengthof materials containing small to moderate critical flaws may be lowerthan that of the pure matrix, even for a composite with a homogeneous


particle distribution, and notwithstanding the fact that the toughness formuch longer cracks can exceed that of the matrix by a considerable margin.

∑ For a given composite system, the reduction in strength compared withthe unreinforced matrix material is often greatest for the composites showingthe greatest amount of steady-state toughening because bridging is mosteffective with large particles owing to the retention of high bridging forcesfor greater crack face separations. Larger particles also imply an increasedparticle spacing for a given volume fraction, however, so cracks mustgrow further before encountering the first bridging element that is responsiblefor the increasing toughness.

∑ Finally, materials with very large flaws, such as may result from heavysurface damage, benefit most from the toughening effect. The weakestspecimens in a batch of particle-toughened composites may, therefore, beconsiderably stronger than those in a batch of pure matrix material, andthe range of strengths present is reduced. This is most important from thepoint of view of applications for these materials because it is the weakestspecimens that determine the usable strength, and the smaller range ofstrengths present and resistance to damage are attractive to designers.

The Al2O3–Al2TiO5 particulate composites investigated by Bennison andco-workers [7, 31] provide an excellent illustration of several of these effects.The volumetric thermal expansion of aluminium titanate, which has anorthorhombic crystal structure, is similar to that of alumina, but unlike aluminait exhibits extreme anisotropy of linear expansion. This leads to sizeablethermal stresses in the composites which are at least partly responsible forthe significant crack interface bridging observed [7]. Figure 4.7 shows alog–log plot of strength against indentation load for biaxial bend tests withVickers hardness indentations at the point of maximum tensile stress. Theindentations are a method of introducing controlled flaws to the specimens,and simulate contact damage. The experimental points on the plot are froman Al2O3–20vol%Al2TiO5 composite with a grain size of 6 mm. The meanstrength of this material when unindented was ~250 MPa. This is considerablybelow the strength of pure alumina with the same grain size, which can beestimated as ~510 MPa by interpolation of strengths obtained for other grainsizes in the same laboratory [32]. The strength of the composite is remarkablyinsensitive to indentation, however, and remains substantially unaltered evenwith indentation loads as large as 300 N. The experimental curve (A) foralumina with a grain size of 2.5 mm and the prediction for an alumina withthe same grain size as the composite, 6 mm (B), are substantially moresensitive to indentation load, and for indentation loads of around 30 N ormore become weaker than the composite, the strength falling to around halfthat of the composite for the highest indentation loads used. Alumina with agrain size of 80 mm (C) is also flaw tolerant, but is significantly weaker than


the composite. Another attraction of this composite system is that it can bepressurelessly sintered in air, either by mixing Al2O3 and Al2TiO5 powders,or by reaction sintering [33].

4.5.2 Ceramic nanocomposites

Figure 4.7 shows that particulates of several micrometres in dimension canyield desirable properties such as flaw tolerance, but there is a limit to howfar this approach can be extended owing to the inverse relationship frequentlyobserved between strength and toughness. An alternative approach has beenpioneered by Niihara and his co-workers, who in 1991 published a review ofresults from a variety of ‘ceramic nanocomposites’ [34], i.e. composites inwhich at least one of the phases is ‘nanoscale’. The most striking results arefrom microstructures consisting of alumina grains of conventional size (afew microns) containing SiC particles with a mean diameter of ~0.25 mm.The addition of 5 vol% submicron SiC to alumina was reported to increaseits strength by a factor of three, from 350 MPa to 1050 MPa. Annealing ofthe composites increased the strength further to 1520 MPa.

These results have proved to be controversial because they have beendifficult to reproduce, although it should be stated at the outset that the highstrength of Niihara’s materials has been verified independently. The mainarea of uncertainty is over the extent of the strength increase on adding SiC

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B

C

4.7 Plot of strength versus indentation load for alumina with variousgrain sizes and alumina/aluminium-titanate composite. The datapoints and solid line are from the composite. The other lines are foralumina with various grain sizes: A, 2.5 mm, experimental; B, 6 mm,interpolated; C, 80 mm, experimental (reproduced from Bennison etal. [7] by kind permission of Taylor and Francis Ltd (http://www.tandf.co.uk/journals).


(in the absence of any annealing treatment). Subsequent investigations byother workers in which the strengths of alumina/SiC nanocomposites havebeen compared with those of alumina processed in the same way and withthe same grain size and grain size distribution (i.e. no abnormal grain growthin either material) also show the nanocomposites to be the stronger material,but typically by only 10–50%, e.g. [13, 35, 36]. Several factors appear tocontribute to this inconsistency. One is that the extent of the strength increaseoriginally reported by Niihara [34] is partly attributable to the unusually lowstrength of the alumina used in the comparison, as well as to the high strengthof the composite. Such low strengths in alumina are normally indicative ofprocessing defects such as porosity or abnormally grown grains, neither ofwhich was present in the nanocomposites. At the other end of the scale,strengths of ~1000 MPa have previously been achieved in monolithic aluminaby using a special processing technique to break down powder agglomerates[37]. Though this processing method was not used in producing thenanocomposites in Niihara’s work [34], it is clear that a full assessment ofthe effect of the SiC nanoparticles on strength can only usefully be madeusing materials that differ only in whether they contain SiC or not. Eventhen, composite results from different sources may differ in the size anddistribution of the SiC particles. Agglomerations of SiC particles maybe capable of acting as critical flaws. The distribution of particles, andconsequently the strength of the composite, will thus depend on fine detailsof the processing that are difficult to reproduce precisely from one laboratoryto another [10].

The variable nature of the strength increases in these particulatenanocomposites has also made it difficult to identify the explanation forthem with certainty, though there is wide agreement that the SiC produces nosignificant toughening [13, 35, 36, 38] and that the explanation therefore lieselsewhere. There is some evidence that crack initiation may be inhibited inthe nanocomposites [38], and several explanations for this have been putforward based on interactions between cracks and the large thermal stresses(~ -2 GPa in the particles [39]) which are one of the most obvious consequencesof the SiC particle additions [40, 41]. Another important factor is undoubtedlythat surface damage and machining stress development during grinding andpolishing differ dramatically between alumina and the nanocomposites, withthe nanocomposites exhibiting significantly less surface cracking and pullout.If the critical defects are machining-induced, this can be expected to explainat least partially the increased strength. The response to surface abrasion isdescribed more fully in Section 4.7.2.

The further increase in strength on annealing reported by Niihara hasbeen reproduced in several investigations [35, 42, 43]. The extent of thestrength increase depends on the surface finish [42], and there is convincingevidence that this is caused by crack healing, aided by a glassy phase resulting


from oxidation of SiC particles near the surface that fills the cracks andbonds the faces together [42]. This has been shown to occur even in atmospheressuch as standard laboratory argon that are nominally inert. The few reportsin which annealing did not increase the strength (e.g. [38]) may be aconsequence of subtle differences in the annealing atmosphere (e.g. loweroxygen partial pressure), or failure from subsurface cracks where oxidationcannot take place. This self-healing effect brings to oxide ceramics a potentialadvantage otherwise available only to silicon-containing non-oxides such asSi3N4.

In conclusion, the strengthening of alumina/SiC ‘nanocomposites’ looksset to remain controversial owing to its capricious nature. Commercially, theroom-temperature strength increase is not sufficiently large for it to repaythe extra cost of processing compared with alumina, since the nanocompositesrequire an inert atmosphere and slightly higher temperatures, even withsintering aids such as yttria, which have been shown to alleviate the inhibitionof sintering caused by the SiC additions. The high-temperature propertiesand the polishing and wear behaviour of the nanocomposites offer muchmore significant improvements, however, which may well be cost effective(see Sections 4.6 and 4.7.2).

4.6 High-temperature strength

One of the main drivers for the application of ceramics is their ability tomaintain their strength at high temperature. In monolithic ceramics withoutpotent toughening, there are several stages of high-temperature behaviour.At moderate temperatures, below the level at which solid-state diffusion orother high temperature mechanisms become significant, standard measurementsof toughness and strength show little temperature dependence, although slowcrack growth may be accelerated considerably, particularly in oxide ceramicswhen crack growth is caused by the interaction of water vapour with thematerial at the crack tip. Similarly, composite systems with a small thermalexpansion mismatch between the phases such as Al2O3–TiC exhibit neitherthermal residual stresses, nor in this case strong toughening mechanisms, sothe toughness is moderate and independent of temperature until newmechanisms operate at high temperature.

In particulate composites exhibiting strong toughening mechanisms suchas crack bridging and stress-induced microcracking, a more marked changein toughness and strength might be expected at moderate temperatures owingto the reduction in the thermal residual stresses locked into the microstructure.This would clearly inhibit stress-induced microcracking and may also reduceboth the number of bridging elements formed and the closure force theyexert if the thermal stress clamps them in place. The more minor directtoughening effect of the fluctuating residual stress field would also be reduced.


SiC–TiB2 particle composites provide a good illustration of this, sincemicrocrack toughening, crack bridging and residual stress toughening are allexpected to operate at room temperature. Jenkins, Salem and Seshadri [44]found that the long-crack (chevron notch) fracture toughness of such compositesfell gradually as the temperature was increased from room temperature to1400∞C, at which temperature the toughness was close to that of monolithicSiC (Fig. 4.8), indicating that the toughening mechanisms operative at roomtemperature had ceased to be effective.

Interestingly, toughness values for the SiC–TiB2 particle composites in[44] derived from SENB tests using blunt notches made with a 300 mmdiamond saw did not exhibit such a marked reduction with temperature.Similarly, McMurtry et al. [6] found that the flexural strengths of similarcomposites was independent of temperature between room temperature and1200ºC. This suggests that the initial portion of the R-curve, which determinesthe strength when failure is from small flaws such as those at the tip of asawn notch or the surface of a flexural strength specimen, is not greatlyinfluenced by the toughening mechanisms mentioned.

At very high temperatures, typically in excess of 1000ºC, the deformationand fracture behaviour of monolithic ceramics becomes complicated by theoperation of new mechanisms such as solid-state diffusion, grain boundarysliding, the activation of dislocation slip systems, the melting of thin grainboundary films, and oxidation. All of these can also occur in particulateceramic composites. One example of such effects is the observation of asharp toughness increase, which is well known to be caused by crack bluntingor healing associated with softening of grain boundary phases, followed bya rapid loss of strength with further temperature increases as the grain boundaryphase loses its strength completely [45]. This has been observed in Si3N4–TiC and Al2O3–TiC composites by Baldoni et al. [46], who point out that itis unlikely that the particulate reinforcement plays an important role. At

0 500 1000 1500Temperature (∞C)

4.5

4

3.5

3

2.5

2

Tou

gh

nes

s (M

Pa

m1/

2 )

4.8 Fracture toughness against testing temperature for a commercialSiC–16vol%TiB2 particle composite tested using chevron notchedbeams in three-point bending. Data from [44].


higher temperatures, creep and the associated cavitation or cracking associatedwith grain boundary sliding can lead to composite failure in much the sameway as for monolithic ceramics.

Although the processes occurring in particulate composites at hightemperatures qualitatively resemble those in monolithic ceramics, there arenevertheless several examples of particulate additions leading to significantproperty improvements. Modifications to the grain boundary structure,associated phases or segregants are often involved. French et al. [47], forinstance, have reported that duplex microstructures consisting of equalproportions by volume of Al2O3 and yttrium aluminium garnet (YAG) exhibitedcreep rates at 1250ºC that were slower than those of ‘pure’ alumina and YAGby factors of 20 and 4 respectively. This was explained by the ability of yttriaadditions to reduce the creep rate of alumina by approximately two orders ofmagnitude. When the alumina in the composite was considered as beingyttria-doped, the composite obeyed the rule of mixtures. Although the lowestcreep rate was obtained from single-phase yttria-doped alumina, the compositemight be preferable in some situations because of the increased microstructuralstability conferred by the duplex structure. The composite suffered negligiblegrain growth during the creep tests, for instance, but the grain size of theyttria-doped alumina increased noticeably. This can be attributed to the greaterdiffusion distance required for grain growth in multiphase structures.

Another notable example of a reduction in creep rate through the additionof second-phase particles concerns ‘nanocomposites’. In alumina–SiCn systems,several investigations have reported significant reductions in creep ratecompared with monolithic alumina [48, 49]. Figure 4.9 shows the results ofOhji and co-workers [48]. At 1200ºC the creep rate of an Al2O3–17vol%SiCnanocomposite was less than that of alumina for a given stress by a factor of250, and the time to rupture at 50 MPa was increased from 120 h to 1120 h.The SiC inhibits creep primarily because it is difficult to remove or deposit

alumina nanocomposite

1.4 1.6 1.8 2 2.2Log stress (MPa)

– 5

– 6

– 7

– 8

– 9

– 10

Log

str

ain

rat

e (s

–1)

4.9 Log–log plot of tensile creep strain rate against applied stress foralumina and an alumina–17vol%SiC nanocomposite tested in tensionat 1200∞C.


material at the interface with the alumina matrix to allow diffusional transportto occur. Intergranular particles therefore inhibit diffusion creep and grainboundary sliding in the same way that they inhibit sintering (Section 4.2)and, through their consequent immobility, prevent grain growth. The smallerimprovement in time to rupture in these observations shows that the strain tofailure was reduced by the SiC additions. This is attributable to the nucleationof cavities at the intergranular SiC particles.

The suppression of creep has also been reported in Si3N4–SiCnanocomposites, and similar explanations have been given [34, 50, 51],although others have found no improvement [52]. The reasons for thesediscrepancies have yet to be resolved, but it is likely that they originate in thedifferent processing methods and sintering aids used in producing thesematerials and hence the differences in grain boundary phases, as well as inthe wide variety of other additive-induced microstructural variations possiblein Si3N4 materials (e.g. the presence of elongated, whisker-like grains).

As well as being used to inhibit creep, second-phase particle additionscan be used under different conditions to achieve the opposite, in fabricatingceramic microstructures that enable superplastic deformation. This term refersto the ability to achieve large, uniform tensile elongations (�100%) atmoderate strain rates (10–5–10–4 s–1) without failure. The underlying mechanismof this type of deformation involves diffusion, and the main requirements arethat a fine grain size (of the order of microns or finer) can be maintained atthe high temperatures necessary to give rapid deformation at sufficiently lowstresses to avoid failure. A common strategy for producing and maintaininga fine grain size is to use microstructures comprising two or more mutuallyinsoluble phases, often in roughly equal volume fractions. This severelylimits grain growth as described in connection with duplex Al2O3–YAGcomposites above, and many superplastic particulate composites have nowbeen reported. An alternative method of maintaining a fine grain size is touse a lower volume fraction of fine second-phase particles which can restrictgrain growth by Zener pinning. These include ZrO2–Al2O3 [52], ZrO2–mullite[54] and Si3N4–SiC [34, 55]. Perhaps the most impressive results to date arethose of Hiraga and co-workers [56], who have fabricated three-phasemicrostructures consisting of zirconia, magnesium aluminate spinel and a-alumina with a mean grain size of around 200 nm which are capable oftensile elongations of >1000% with a strain rate of 0.4 s–1 at a temperatureof 1650∞C. These exceptional results stem from the combination of the hightesting temperature and very fine and stable grain size that the three-phasestructure allows.

Superplastic ceramics have several obvious potential advantages forcommercial application. These include net size and shape forming and thepossibility of forming complex components from initially flat sheets. Whilstthe practical problems of forming at temperatures in excess of 1200∞C obviously


add cost to the process, the diamond machining which is the only practicalcompetitor for the production of many complex shapes to high dimensionaltolerance is also expensive. Despite these attractions, the phenomenon remainsa scientific curiosity at the time of writing. This is much the same as thesituation for superplastic metals until the late 1960s, when a few practicaldemonstrations of their commercial benefits led to their widespread application.It remains to be seen whether the industrial superplastic forming of ceramicswill take off in the same way.

4.7 Wear

4.7.1 Microcomposites

Another primary motivation for the use of ceramics in engineering applicationsis their high wear resistance. At its simplest, wear involves plastic deformation-controlled mechanisms such as cutting or ploughing and, in ceramics, theremoval of pieces of material by brittle fracture (‘pullout’). This is the originof figures of merit for wear of the form K Hm n

c , where Kc is the toughnessand H the hardness, and m and n are positive exponents. In reality, however,these wear mechanisms are much more complex than this suggests, with theformation of modified surface microstructures and compacted layers beingcommon, and additional mechanisms such as chemical interaction betweenceramic and substrate, or atmosphere and ceramic, are frequently important.The high temperatures generated locally during the wear process add to thiscomplexity. Even in cases where it can be argued that the simple plasticityor brittle fracture mechanisms are dominant, the appropriate values of Kc

and H to use in models are not clear, as the scale of the plastic deformationor fracture is much smaller than that in tests used for the measurement ofthese properties, and the temperature at which these properties should bemeasured is ill defined. Furthermore, the dominant mechanism and the rateat which it operates depend not only on the ceramic itself, but on the wearconditions and substrates involved.

Many of the reports of wear tests on particulate ceramic composites areabrasive tests (e.g. grinding on different grades of SiC paper [57]) ormeasurements associated with specific applications, the outstanding examplebeing cutting tools, in which this class of composite finds widespreadapplication. The agreement in raw results from different studies is sometimescontradictory. Sarin et al. [58, 59], for instance, tested a range of compositesby abrasion with dry 45 mm diamond in argon, and found that the wearresistance scaled approximately with K Hc

3/4 1/2 and the order of ranking ofthe materials tested (lowest wear resistance first) was Al2O3, Al2O3–ZrO2,Al2O3–TiC, Si3N4+Y2O3, SiAlON and Si3N4–TiC composite. Holz et al.,however, performed abrasive pin-on-disc tests on a wide range of composites


using two grades of SiC paper with 15 mm and 70 mm grit particles, andfound no clear correlation with a relationship of the form K Hm n

c [57]. Theaddition of ZrO2 to alumina had little effect its wear resistance in this study,but the further addition of TiC/TiN reduced the wear rate by a factor of ~3to produce one of the most wear-resistant materials tested, which, along witha hot-pressed monolithic b-SiAlON, was far superior to either of the twoSi3N4–TiC/TiN composites tested.

The sensitivity of wear to so many experimental factors is undoubtedly amajor part of the reason for some of these apparently contradictory conclusions.Another is that important details of the microstructures of the materialsbeing compared, such as the matrix grain size, particle size and amount ofporosity, differ between the two studies. Such features can have a profoundeffect on the wear rate. Indeed, although the original motivation for addingTiC particulate to Al2O3 cutting tools was that TiC was harder, stiffer andmore thermally conductive than alumina, though difficult to process as amonolith, it is now thought that the main reason why the particulate improvesthe hardness, strength and wear resistance is its grain refining effect [4].

The success of Al2O3–TiC cutting tools for machining steels and cast ironis interesting in the context of the good bonding between the particles andthe matrix, the small thermal expansion mismatch, and consequently thelimited amount of toughening in this composite system [60]. In toughercomposites, the microstructural features such as thermal stresses and weakinterfaces which are instrumental in the operation of toughening mechanismssuch as crack bridging, microcracking, crack deflection and the directtoughening effect of residual stresses are also a potential aid to the initiationand propagation of the short, near-surface cracks that are responsible forsevere wear by surface fracture and pullout, and so are potentially damagingto the wear resistance. Holz et al. [57], for instance, observed that SiCplatelets in a reaction-bonded silicon nitride matrix were only weakly bondedto the matrix and pulled out during abrasive wear, increasing the wear rateboth directly, and indirectly by acting as abrasive particles themselves.

In cases where the mechanical properties of the composites play a dominantrole, the extent to which the ease of crack initiation in toughened compositesis overcome by the potential advantages of the particles (higher hardness,global toughness, etc.) depends on the relative scales of the microstructureand the cracks caused by abrasion. If the microstructural scale is greater thanthe crack size, the wear resistance is likely to be degraded, and vice versa.Figure 4.10 shows the effect of microstructural scale on the abrasive wearresistance of Si3N4/TiCp composites from the work of Wayne and Buljan[61]. Fine-scale composites, with a mean TiC particle size of 0.4 mm, improvethe wear resistance relative to monolithic Si3N4, but coarser particles (≥ 1.5mm) degrade it. The same work also demonstrates the influence of the typeof test on the ranking of materials: composites with all three particle sizes


showed significant wear rate reductions relative to monolithic Si3N4 whensubjected to a gas-jet particle erosive wear test.

In other cases, the effect of particulate additions is dominated by chemicalrather than mechanical effects, usually in less aggressive tests or in reactiveenvironments. There is now a substantial body of literature describing thebeneficial effects on wear rates of adding titanium-containing particles toceramic matrices. These form a soft adherent tribochemical film of Ti-containing oxide that reduces the wear rate of the underlying surface fromdamage. Examples include unlubricated ball (Si3N4 or steel)-on-disc wear ofSi3N4–TiB2 in laboratory air [62] and SiC–TiC and SiC–TiC–TiB2 in oscillatingsliding against SiC and a-Al2O3 in water [63].

4.7.2 Nanocomposites

In view of the above comments on the importance to wear resistance of afine microstructural scale, it is perhaps not surprising that the most remarkableeffect of submicron particle additions in alumina/SiC ‘nanocomposites’ is toincrease the resistance to severe wear dramatically compared with monolithicalumina. Zhao et al. [35] first commented that the nanocomposites wereeasier to polish to a mirror finish than alumina, and significant improvementsin surface finish are also found for more aggressive surface treatments asFig. 4.11 demonstrates [9]. Figure 4.11(a) shows the surface of a piece ofalumina ground with 45 mm diamond paste. Grain facets resembling aconventional fracture surface can be seen behind the wear debris, showingthat much of the material has been removed as large pieces by intergranularfracture around their periphery. Figure 4.11(b) shows an alumina–11vol%SiCnanocomposite, which has the same alumina grain size, after being subjectedto the same treatment. The appearance of the surface contrasts sharply with

0 1 2 3 4Particle size (mm)

Si3N4

6

5

4

3

2

1/V

(10

5 cm

–2)

4.10 Wear resistance, expressed as the reciprocal of the abradedvolume, V, against microstructural scale, represented by the TiCmean particle diameter, for Si3N3–20vol%TiCp composites. Data fromthe work of Wayne and Buljan [61].


that of the alumina, there being little brittle fracture and pullout, wear havingbecome dominated by plasticity-controlled mechanisms as shown by thescratches covering most of the surface. The addition of only 2 vol% SiCreduces the area fraction of pullout by more than a factor of 2, and thecorresponding reduction with 10% SiC is a factor of more than 50 [64].

Walker et al. [11] were the first to report improved wear resistance, findingthat the addition of SiC reduced the wet erosive wear rate by a factor of 3compared with alumina with the same grain size. It has now become clearthat this is a quite general observation in conditions of severe wear, withreductions in wear rate of a similar order having been found for abrasivewear [64] and dry sliding wear [65]. In addition, Chen et al. [66] have shownthat the mild to severe wear transition seen in sliding wear of alumina iseither delayed or completely suppressed in the nanocomposites. They alsofound that there is no improvement in wear resistance in mild wear, wheresurface fracture is absent.

Recent work correlating microstructure, abrasive wear rate and theappearance of the worn surfaces [64] has concluded that the most importantreason for the improved surface finish and wear resistance on adding SiC isthat the dimensions (diameter, depth) of the individual pullouts are reduced.This was explained in terms of the well-established change in fracture modefrom intergranular in alumina to transgranular in the nanocomposites, whichallows pullouts smaller than the grain size to be formed when fracture initiatesin the nanocomposites, whereas in alumina, near-surface cracks tend to followgrain boundaries, giving a minimum pullout dimension of the order of thegrain size. SiC additions of 10 vol% were also shown to inhibit the initiationof the cracking responsible for pullouts directly, and there is evidence thatthis is because the SiC particles inhibit the subsurface twinning to which

4.11 Comparison of the response of pure alumina and an alumina/SiC nanocomposite to abrasion with 45 mm diamond paste [9]: (a)alumina, grain size 2.6 mm; (b) alumina–11 vol% SiC, alumina grainsize 2.6 mm.

(a) (b)


alumina is prone [67]. It is also worth noting that the suppression of fractureinitiation and surface damage in the nanocomposites provides a naturalexplanation for at least some of the strength improvement in these compositesdescribed in Section 4.5.2.

4.8 Future trends

The understanding of the mechanical behaviour of particulate composites iswell advanced, and it seems unlikely that step changes in properties will beforthcoming through the discovery of new strengthening mechanisms, atleast for materials with microstructural scales of 100 nm upwards. The mostobvious area for further investigation is in particulate composites based onlength scales smaller than this. Whilst the ‘nanocomposites’ described in thischapter have proved controversial, the striking results shown in Fig. 4.11provide convincing evidence that the continuing refinement of microstructurecan lead to novel and striking effects. The first task to be undertaken is tofind processing methods capable of producing such microstructures, andhere it is likely that research on particulate composites will proceed in closerelationship with the efforts to produce single-phase ceramics with trulynanoscale grain sizes that are beginning to bear fruit [68].

Another area for novel work may be in functional particulate composites.Most functional ceramics in use at present are essentially single phase. Theinsertion of particles into the structure may interact in novel ways with thefunctional elements of microstructure. A simple example might be the pinningof ferroelectric domain wall motion during formation or operation, whichmay improve dielectric and electromechanical losses. Strength improvementsmight also be obtained. A further possible area of interest is the interactionbetween thermal stresses around particles with the strains occurring duringcooling through the Curie temperature of BaTiO3 and other perovskites,which offers the possibility of producing new domain structures.

At present the only commercial use of particulate ceramics of any note isalumina–TiC cutting tools, which as noted above rely more on the grain-refining effect of the particulate rather than on any of the true compositeeffects described. Thus a further area for development is in the application ofthese ceramics. The lack of application to date should not come as a surprise.The history of new materials has often shown a lag between development inthe laboratory and commercial use. Monolithic ceramics are no exception tothis. Despite the remarkable developments in structural ceramics that havetaken place since the 1960s, truly structural applications remain rare. This isbeginning to change, although it is important to realise that the applicationsfor which structural ceramics are well suited are limited in scope by theirnature. Monolithic nitride ceramics are beginning to be used in semi-structuralapplications in jet engines, and it is reasonable to suppose that as confidence


and experience are gained with these, the use of ceramics will spread. Theevidence in this chapter suggests that the potential role of most particulateceramics is to give modest property improvements over monolithic ceramicswith little cost penalty. The last point is important, and realistically requiresthe use of particles that allow pressureless sintering and that are chemicallycompatible with the matrix in terms of processing environment and oxidation.For improving the flaw tolerance of alumina, for instance, Cr3C2 particlesare unlikely to be cost effective since they require the use of an inert orreducing environment. Al2TiO5 particles, however, allow pressureless sinteringin air and are therefore an excellent candidate for exploitation in aluminacomponents that require tolerance to surface damage. The alumina/SiC‘nanocomposites’ may be an exception to this rule, in that the improvementsin wear resistance may be sufficiently spectacular for it to be worth payingthe premium associated with sintering in inert gas. Instead of regarding thesematerials as a superior (but expensive) alternative to alumina, it may bebeneficial to market them as a cheap alternative to SiC or Si3N4.

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35. Zhao, J., Stearns, L.C., Harmer, M.P., Chan, H.M., Miller, G.A. and Cook, R.E.‘Mechanical behaviour of alumina–silicon carbide “nanocomposites” ’, J. Am. Ceram.Soc. 76 (1993) 503–510.

36. Pérez-Rigueiro, J., Pastor, J.Y., Llorca, J., Elices, M., Miranzo, P. and Moya, J.S.‘Revisiting the mechanical behaviour of alumina/silicon carbide nanocomposites’,Acta Mater. 46 (1998) 5399–5411.

37. Alford, N.M., Birchall, J.D. and Kendall, K. ‘High-strength ceramics through colloidalcontrol to remove defects’, Nature 330 (6143) (1987) 51–53.

38. Meschke, F., Alves-Riccardo, P., Schneider, G.A. and Claussen, N. ‘Failure behaviorof alumina and alumina/silicon carbide composites with natural and artificial flaws’,J. Mat. Res. 12 (1997) 3307–3315.

39. Todd, R.I., Bourke, M.A.M., Borsa, C.E. and Brook, R.J. ‘Neutron diffractionmeasurements of residual stresses in alumina/SiC nanocomposites’, Acta Mater. 45(1997) 1791–1800.

40. Jiao, S. and Jenkins, M.L. ‘A quantitative analysis of crack-interface interactions onalumina-based nanocomposites’, Phil. Mag. A78 (1998) 507–522.

41. Hoffman, M. and Rödel, J. ‘Suggestion for mechanism of strengthening of“nanotoughened” ceramics’, J. Ceram. Soc. Japan 105 (1997) 1086–1090.

42. Wu, H.Z., Lawrence, C.W., Roberts, S.G. and Derby, B. ‘The strength of Al2O3/SiCnanocomposites after grinding and annealing’, Acta Mater. 46 (1998) 3839–3848.

43. Ando, K., Chu, M.C., Tsuji, K., Hirasawa, T., Kobayashi, Y. and Sato, S. ‘Crackhealing behviour and high-temperature strength of mullite/SiC composite ceramics’,J. Eur. Ceram. Soc. 22 (2002) 1313–1319.

44. Jenkins, M.G., Salem, J.A. and Seshadri, S.G. ‘Fracture of a Tib2 particle/SiC matrixcomposite at elevated temperature’, J. Comp. Mat. 23 (1989) 77–90.

45. Davidge, R.W. ‘Mechanical Behaviour of Ceramics’, Cambridge University Press,Cambridge, UK, 1979.

46. Baldoni, J.G., Buljan, S.T. and Sarin, V.K. ‘Particulate titanium carbide–ceramicmatrix composites’, Proc. 2nd Int. Conf. Science of Hard Materials, Inst. Phys.Conf. Ser. 75, IoP/Adam Hilger, Bristol, UK, 1986.

47. French, J.D., Zhao, J. Harmer, M.P. Chan, H.M. and Miller, G.A. ‘Creep of duplexmicrostructures’, J. Am. Ceram. Soc. 77 (1994) 2857–2865.

48. Ohji, T., Nakahira, A., Hirano, T. and Niihara, K. ‘Tensile creep behavior of alumina/SiC nanocomposite’, J. Am. Ceram. Soc. 77 (1994) 3259–3262.

49. Thompson, A.M., Chan, H.M. and Harmer, M.P. ‘Tensile creep of alumina–siliconcarbide “nanocomposites” ’, J. Am. Ceram. Soc. 80 (1997) 2221–2228.

50. Niihara, K., Suganuma, K., Nakahira, A. and Izaki, K. ‘Interfaces in Si3N4–SiCnanocomposite’, J. Mat. Sci. Lett. 9 (1990) 598–599.

51. Park, H., Kim, H.E. and Niihara, K. ‘Microstructure and high-temperature strengthof Si3N4–SiC nanocomposite’, J. Eur. Ceram. Soc. 18 (1998) 907–914.

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54. Yoon, C.K. and Chen, I.W. ‘Superplasticity of two-phase ceramics containing inclusions– zirconia mullite composites’, J. Am. Ceram. Soc. 73 (1990) 1555–1565.

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Part II

Graded and layered composites

129

131

5.1 Introduction

Layered-graded materials (LGMs) exhibit a stepwise or progressive changein composition, structure, and properties as a function of position within thematerial [Koizumi, 1993; Hirai, 1996]. This innovative design eliminatesever-present sharp boundaries in conventional composites which may impartundesirable physical and mechanical properties. An example is debonding orseparation at the boundary due to thermal or residual stress induced bymismatch in thermal expansion. LGMs have been processed by a variety ofmethods. A list of the common synthesis methods & examples of thesematerials have been described elsewhere [Sakai & Hirai, 1991]. Recently,liquid infiltration of preforms [Marple & Green, 1990; Low, 1998a] hasemerged as an innovative technique for the processing of graded compositematerials. Using this infiltration process, it is possible to design new materialswith unique microstructures (e.g. graded, multiphase, microporous, etc.) andunique thermomechanical properties (e.g. graded functions, designed residualstrains, thermal shock, etc.).

Recent developments in layered ceramics have provided a strategy forlaminating the ceramic structure with an outermost homogeneous layer toprovide wear resistance and an underlying heterogeneous layer to providetoughness [An et al., 1996; Liu et al., 1996; Padture et al., 1995]. Theselayered structures promote toughness by interlayer crack deflection throughweak interfacial bonding, or strength by incorporating macroscopic compressiveresidual stresses through strong interlayer bonding. Layered ceramics producedin this second way have shown uncommonly high damage resistance underHertzian loading, with retention of strength and wear resistance. However,these layered structures are disadvantaged by either the counterproductiveeffects of weak interlayers or the excessively large residual stresses that cancause enhancement of delamination failures.

Here we consider a new approach, in which microstructural elements aretailored to provide graded compositions and generate different modes of

5Functionally-graded ceramic composites

I M L O W , R D S K A L A and P M A N U R U N G,Curtin University of Technology, Australia


strengthening and toughening. The basic idea is to produce a graded dispersionof particles within the alumina matrix through an infiltration process to yielda layer of homogeneous alumina for hardness and wear resistance, and aheterogeneous layer of tough graded alumina for damage dispersion. Thedesign of this layered-graded material can provide a unique mechanicalperformance of both flaw tolerance and wear resistance. The proposed strategyof designing layer structures with graded interfaces for crack arrest isfundamentally different from the conventional laminar approach, where thereis a sudden change in composition at the interface between layers. Thisabrupt interface can cause cracking or delamination due to thermal expansionand elastic modulus mismatches. The toughening processes envisaged, inwhich the heterogeneous graded layers act to inhibit crack penetration byinterlayer ‘stress shielding’ and by ‘crack bridging’, offer major advantagesover conventional layer composites, where toughness is introduced viadeflection of transverse cracks along weak interfaces.

In this chapter, we describe the synthesis and characterisation of themicrostructure and properties of layered-graded alumina-matrix compositesthrough liquid infiltration. This approach is relatively simple and offersexcellent control over the depth of the graded layer. The presence of a gradeddispersion of reinforced particles in the alumina matrix has a profound influenceon the physical and mechanical properties of the composites. An overview ofthe infiltration kinetics and the use of the infiltration process as a new philosophyfor tailoring novel graded ceramic systems are also presented.

5.2 Infiltration kinetics and characteristics

Liquid-phase infiltration of preforms has emerged as an extremely usefulmethod for the processing of composite materials. This process involves theuse of low-viscosity liquids such as sols, metal- or polymer-melts. Using thisinfiltration process, it is possible to design new materials with uniquemicrostructures (e.g. graded, multiphase, microporous) and uniquethermomechanical properties (graded functions, designed residual strainsand thermal shock).

Liquid infiltration into dry porous materials occurs due to capillary action.The mechanism of infiltrating liquids into porous bodies has been studied bymany researches in the fields of soil physics, chemistry, powder technologyand powder metallurgy [Carman, 1956; Semlak & Rhines, 1958]. However,the processes and kinetics of liquid infiltration into a powdered preform arerather complex and have not been completely understood. Based on Darcy’sfundamental principle and the Kozeny–Carman equation, Semlak & Rhines(1958) and Yokota et al. (1980) have developed infiltration rate equationsfor porous glass and metal bodies. These rate equations can be used todescribe the kinetics of liquid infiltration in porous ceramics preforms, but

Functionally-graded ceramic composites 133

a single capillary model will result in the pore size measured from theexperimental infiltration rates being an order of magnitude smaller than thepore sizes seen from SEM and porosimetry measurements [Dullien et al.,1977; Einset, 1996]. Low and co-workers (2000) have shown that a combinationof the Washburn model [Washburn, 1921] and Dullien’s analysis [Dullien etal., 1977] is able to reconcile the infiltration rates and the pore sizes determinedfrom SEM and porosimetry measurements for the infiltration of water andTiCl4 into alumina preforms.

5.2.1 Modelling infiltration kinetics

There are a number of formulae which are relevant for modelling the infiltrationkinetics of a liquid into preforms. The first equation to calculate the heightof infiltration against time was formulated by Washburn (1921):

hr

t = cos 2

1/21/2g q

hÊË

ˆ¯ (5.1)

where h and t are the height of liquid and the time, respectively, and g, q, rand h are the surface tension, contact angle, pore radius and viscosity,respectively. Hence, the rate of infiltration (i.e. h/t) is a function of surfacetension, contact angle, pore radius and viscosity.

Another model, proposed by Yokota et al. (1980), involves tortuosity, T,and shape factors, Cs:

hT

C p rt = 2

cos s1/2

1/2g qh

ÊË

ˆ¯ (5.2)

The value for T is normally 1.4142 and Cs is 0.4. When water wets thepreform completely and liquid is spreading, the contact angle q is 0∞ [Yokotaet al., 1980]. A third formula was proposed by Travitzky & Shlayen (1998):

h rP

t = 2

net1/2

1/2

hÊË

ˆ¯ (5.3)

where the net pressure, Pnet is given by:

Pr

gh Pnet atm = 2 cos

– – g q r (5.4)

Here r is the density of the infiltrant, g is the acceleration due to gravity (9.8m/s2) and Patm is the pressure of the surrounding space above the liquidinfiltrant. This model is useful for predicting the influence of pressure on therate of infiltration. Another formula is the Kozeny–Carman equation [Carman,1956]:


h ht

kT

dd

= (5.5)

where T is the tortuosity of the capillary tube and k is a constant.

5.2.2 Studies on infiltration kinetics

Mercury porosimetry measurements for a partially sintered alumina preformshowed a bimodal pore size distribution with neck diameter Dn = 0.15 mm[Manurung, 2001]. As a comparison with the pore sizes and distribution ofthe preform measured by porosimetry, SEM micrographs (Fig. 5.1) weretaken before and after infiltration. Based on SEM examination, the pores inthe preform before infiltration ranged in size from r ~ 0.1–0.5 mm. Assumingan average pore radius of 0.3 mm, this radius is approximately four timeslarger than the pore-neck radius (Dn = 0.15 mm, so pore radius = 0.075 mm)determined by mercury porosimetry.

It is interesting to note that the pore sizes appear to remain virtuallyunchanged following the infiltration process. This suggests that (i) the capillaryforces involved did not appear to cause any shrinkage or size reduction ofthe pores, and (ii) the infiltrant had not filled up the pores but only formeda thin layer deposit on the walls of pores. The small particles entering thegrains are believed to arise from the infiltrant which had entered the porechannels and adhered to the walls of the pores following drying [Manurung,2001].

In order to successfully model the infiltration kinetics in terms of theeffects of presintering temperature, type of infiltrant, infiltration environment,and multiple infiltrations, the pore radius of alumina preform (presintered at1000∞C) was measured using water as infiltrant, since the viscosity and

(a) (b)

5.1 SEM micrograph of the partially sintered alumina preform(1000∞C) showing the pore microstructure (a) before and (b) afterinfiltration with TiCl4 [Manurung, 2001].


surface tension of water are well known. The pore radius of the aluminapreform was calculated from the Washburn model.

By assuming a contact angle q = 0∞ [Einset, 1996; Ligenza & Bernstein,1951] the pore radius of the preform can be calculated if the height and timeof infiltration are known. The rate of infiltration is determined from theslope in Fig. 5.2(a) and then from this slope the pore radius can be found.From the measurements, it was found that the pore radius of the aluminapreform is 0.015 ± 0.001 mm. Similarly the pore radius found from aluminapreform infiltrated with TiCl4 is 0.018 ± 0.002 mm [Manurung, 2001]. Theerrors in the radii only reflect the experimental uncertainty in the measuredvalues for surface tension and viscosity. However, the measured pore radiusis an order of magnitude smaller than the pore radius determined fromporosimetry and SEM (Fig. 5.1).

The pore radius determined from infiltration kinetics can be reconciled tothe experimentally determined radii from SEM and porosimetry, by assuminga two-pore-size model (pore neck and pore bulge), instead of a single capillarypore-size [Dullien et al., 1977; Einset, 1996]. The schematic diagram for thistwo-pore-size model is shown in Fig. 5.2(b).

Dullien (1979) considered the rate of capillary rise of a fluid in a modelthree-dimensional network pore structure consisting of a repeating pore elementwith step changes in diameter. The effective diameter, Deff, model is givenby:

D D DDDk

kk

kj

k

jeff

2 3 –1

= 13

S S SÈÎÍ

˘˚

ÊËÁ

ˆ¯

È

ÎÍÍ

˘

˚˙˙

(5.6)

Db

Dn

Infiltrationdirection

0 20 40 60 80Square root of time (s1/2)

(a)

0.07

0.06

0.05

0.04

0.03

0.02

0.01

0

Hei

gh

t (m

)

(b)

5.2 (a) Height of infiltration of water into alumina preforms (sinteredat 1000ºC) as a function of square root of time; (b) a schematic of thetwo-size single-capillary. Db is the pore-bulge diameter, Dn is thepore-neck diameter.


where the summations are over the number of segments of the repeating poreunit.

The effective pore diameters are therefore smaller than the individualdiameters of the pore segments. Consider a two-pore-size repeating unitmodel in equation (5.6), i.e. j = 1, 2 and k = 1, 2 for pore necks and porebulges as shown in Fig. 5.2(b). Thus Db can be calculated by substituting Deff

and Dn into equation (5.6). The results are 0.27(3) and 0.18(4) mm for waterand TiCl4 respectively. The calculated value of Db for both infiltrants is inreasonable agreement with the average pore diameter estimated from SEM(Fig. 5.1). The kinetics of liquid infiltration in porous alumina preforms havebeen found to depend on several parameters such as surface tension andviscosity of infiltrant, porosity and pore size of preforms, and pressure[Manurung, 2001]. An enhanced infiltration rate is most favourable when (a)a preform has a high porosity (>45%), (b) an infiltrant has a low viscosity,(c) the infiltrate is in vacuum due to a greater driving force, and/or (d) thereare multiple infiltrations by virtue of a self-lubrication effect [Manurung,2001].

A more likely reason for the rate increase during subsequent cycles can beattributed to a decrease in the tortuosity. The path that the infiltrant takesduring each cycle is smoothed by the preceding infiltration cycle. Thereforethe tortuosity is decreased. Inspection of equation (5.2) shows that theinfiltration rate will increase if the tortuosity is decreased. From cycle 1 tocycle 3, the rate increases by a factor of approximately 2 (Fig. 5.3). Thetortuosity would have to decrease by a factor of ~2 if the rate increase wasattributed entirely to a decrease in tortuosity. This is unlikely, therefore it issuggested that there may be a slight increase in the pore radius and shapefactor during subsequent cycling, as well as a decrease in tortuosity. It issuggested that the major cause in the rate increase in going from cycle 1 tocycle 3 can be attributed to a decrease in the tortuosity.

The experimental results show that the kinetics of infiltrating water andtitanium tetrachloride into an alumina preform are parabolic with time (Fig.5.3). It has been shown that the viscosity of infiltrants influences the rate ofinfiltration and that the rate of infiltration is pressure dependent. Theexperimental result of faster kinetics in a vacuum as opposed to that in 1 atmagrees with Travitzky & Shlaken (1998) model. The presintering temperaturehas a strong influence on the kinetics of TiCl4 in alumina preforms andmultiple infiltrations increase the rate of infiltration. According to Yokota etal. (1980) the increase in the infiltration rate due to multiple cycling ispredominantly attributable to a decrease in the tortuosity of the preformduring subsequent cycles.


5.3 Infiltration processing of LGMs

Two steps are involved in the design of LGMs using the infiltration process.The first involves the fabrication of a partially sintered porous preformwhich is then followed by its infiltration with an appropriate infiltrant suchas TiCl4, Si(OC2H5)4 or calcium acetate solution. The infiltrated preform issubsequently heat-treated at elevated temperatures to form the desired phasesin situ. The process of infiltration can involve either partial or completeimmersion of the preform in the infiltration. The latter gives rise to sampleswith an outer graded layer and an inner core of the host material [Marple &Green, 1993]. In contrast, the former produces an outer layer of host materialand an inner layer with a graded composition [Low, 1998a]. The depth orthickness of the graded layer can be further increased by multiple cycles of

Water

TiCl4

0 20 40 60 80 100 120Square root of time (s1/2)

(a)

0.06

0.05

0.04

0.03

0.02

0.01

0

Hei

gh

t (m

)

Cycle 3 Cycle 2 Cycle 1

0 10 20 30 40 50 60 70 80 90Square root of time (s1/2)

(b)

0.06

0.05

0.04

0.03

0.02

0.01

0

Hei

gh

t (m

)

5.3 (a) Effect of infiltrants using water and titanium tetrachloride(TiCl4); (b) effect of multiple infiltrations with TiCl4 under vacuum.


infiltration. A typical infiltration process for the alumina/AT and other systemsis shown in Fig. 5.4.

5.4 Characterisation and properties of alumina-

matrix LGMs

5.4.1 Alumina/mullite and mullite/ZTA/mullite systems

Electron microprobe analysis of concentration profiles across sections of thesintered samples revealed the existence of concentration gradients, the mullitecontent decreasing with increasing distance from the surface of the bodies[Marple & Green, 1993]. SEM examination also revealed a microstructuraleffect: the alumina grain size tended to increase from the surface of the

Pellet/bar

Presintering1000–1200∞C [2 h]

Preform(~ 40–45%)

Infiltration in TiCl4solution

Sample

TiCl4solution

TiCl4solution

Functionally-A/AT

Heat1400∞C [12h]1650∞C [2h]

Sample

Fullinfiltration

Partialinfiltration

5.4 Liquid infiltration processing of layered-graded alumina-matrixcomposites.


samples. This suggests that the presence of mullite limits grain growth inalumina.

From evaluating the mechanical properties of the resulting ceramiccomposite material, large increases, of up to 60%, for both the strength(biaxial flexure) and indentation fracture toughness were achieved (see Table5.1). These increases were attributed to the presence of the mullite case andthe resulting residual compressive surface stresses due to the thermal expansionmismatch between the mullite and alumina [Marple & Green, 1991]. Thestrength of the indented samples deviated from the ideal behaviour, indicatingthe possibility of R-curve behaviour in these materials.

The abundance of mullite in the mullite/ZTA system increased withincreasing infiltration time [Low et al., 1993]. The density (r) and mullitecontent of the sintered sample as a function of infiltration time are shown inTable 5.2. The results suggest that the infiltration process was time (t) dependentand diffusion-controlled with the infiltration front travelled as a function oft1/2. The content of mullite was greatest near the surface and decreasedsharply towards the core of the sample.

The presence of mullite and hence compressive surface stresses appearsto improve the hardness and fracture toughness (see Table 5.2). These valuesare at least two to three times higher than those reported for the mullite/alumina system described above. Clearly, the presence of mullite is desirablefor inducing compressive stresses in the vicinity of the surface region byvirtue of the mismatch in thermal expansion between ZTA and mullite. Thissignificant improvement in the observed fracture toughness was attributed to

Table 5.1 Mechanical properties of graded mullite/alumina

Mullite content Strength Modulus Toughness(vol%) (MPa) (GPa) (MPa.m1/2)

0 305 400 3.95 400 375 5.5

10 425 370 5.513 440 355 6.519 510 335 7.0

Table 5.2 Physical and mechanical properties of graded mullite-ZTA

Infiltration Density Mullite Hardness Toughnesstime (h) (g/cm3) content (GPa) (MPa.m1/2)

(vol%)

0 4.15 0 1679 8.01 4.10 4.7 1692 10.54 4.00 5.9 1693 11.56 4.06 6.2 1733 13.0


the presence of the compressive stresses near the surface in addition to otherwell-established energy-dissipative processes such as transformationtoughening, microcracking and/or crack deflection [Low et al., 1994]. Thisexplanation concurs with that proposed by Marple & Green [1992] for theirmullite/alumina composites where the presence of residual compressive stresseswas observed to be the major contributor to increases in strength and fracturetoughness [Marple & Green, 1992].

5.4.2 Alumina/aluminium titanate and alumina–zirconia/aluminium titanate systems

Layered-graded Al2O3/AT and Al2O3–ZrO2/AT systems were synthesisedusing TiCl4 or Ti(OC2H5)4 as an infiltrant [Skala, 2000; Low et al., 1996a].Figure 5.5 shows a typical graded microstructure of this material where thecontent of AT is most abundant near the surface and decreases with increasingdepth. SEM micrographs of the cross-section of the graded Al2O3/AT atvarious depths are presented in Fig. 5.6. The alumina grains are a dark colourwhile the AT grains are a lighter shade of grey. Figure 5.6(a) shows themicrostructure in the near-surface regions of the LGM. The presence of ATcan be clearly seen to have a beneficial effect on the growth of the aluminagrains. The alumina grains within the surface regions of the LGM are relativelysmall and equiaxed when compared to those of the core (Fig. 5.6(c)). Thealumina grains within the core region are extremely large and elongated,with some of the grains containing intragranular porosity due to exaggerated

5.5 Back-scattered scanning electron micrograph showing thegradation of AT distribution within the sample. Direction ofinfiltration is from right to left.


grain growth and grain boundary migration. The grain growth behaviour ofthe alumina phase is clearly modified by the presence of AT, probably due toa solid-solution effect. Very similar observations in compositional andmicrostructural gradations were obtained for the layered-graded ZTA/ATsystem [Pratapa, 1997].

LGMs of AT/Al2O3 have also displayed some very unique but interestingproperties which include excellent machinability, low thermal expansioncoefficient, improved thermal shock resistance, low hardness, low Young’smodulus and enhanced tolerance to damage [Low, 1998a, 1998b; Skala,2000; Low et al., 1996a].

The XRD depth profile shows that the top surface region is very rich inAT (~88 wt%) with the concentration decreasing slowly as the depth increasestowards the middle region of the sample. The amount of AT formed here isvery large when compared to that found by other workers (Fig. 5.7). Forinstance, Pratapa (1997) found 46 wt% AT phase on the surface of a gradedAT/ZrO2–alumina system, with the value decreasing substantially to 7 wt%at a depth of 0.8 mm. Low (1998b) found 50 wt% AT on the surface ofgraded AAT composites, decreasing to 10 wt% at a depth of 0.5 mm. SimilarlySkala (2000) obtained 52 wt% on the surface which decreased to only 3 wt%at a depth of 0.5 mm in AAT composite. This suggests that unidirectional anddouble infiltrations with the aid of a plastic shield provide an effectivemethod for increasing the content of AT near the surface and for preventing

(a)(b)

10 mm 10 mm

(c)

10 mm

5.6 Typical microstructures of the graded alumina/AT system at (a)surface, (b) 1.0 mm, and (c) 2.0 mm below surface.


a very rapid decrease in the graded AT content. The existence of compositiongradation at both the nanometre and micrometre scale has been verified bygrazing-incidence synchrotron radiation diffraction (Fig. 5.8) [Singh et al.,2002].

The development of induced residual strains within the alumina layer ofAl2O3-Al2TiO5 bilayers with and without graded interfaces in the temperaturerange 20–1500∞C has been observed from the time-of-flight neutron diffractionin terms of line-shift of the (113) reflection. As would be expected, thepresence of a sharp interface in the non-graded sample resulted in the formationof much higher residual strains when compared to the bilayer sample with300 mm thin graded interfaces, due to mismatch in thermal expansion

This workSkala (2000)Pratapa (1997)Low (1998a)

0 0.5 1 1.5 2 2.5Depth (mm)

100

90

80

70

60

50

40

30

20

10

0

Wei

gh

t p

erce

nta

ge

of

AT

(w

t%)

5.7 Depth profiles for the alumina–AT LGM as obtained by variousresearchers, including this study. Error bars indicate two errorestimated deviations (s).

0 1 2 3 4 5 6Grazing-incidence angle (∞)

0.8

0.7

0.6

0.5

0.4

0.3

0.2

0.1

0

Inte

nsi

ty r

atio

(A

T/a

lum

ina)

5.8 Variation of peak-intensity ratio between AT (110) and thealumina (104) as a function of grazing-incidence angle.


coefficients and elastic moduli between Al2O3 and Al2TiO5 in the former.The presence of graded interfaces serves as a ‘buffer region’ to modulate thedifferences in material properties within the bilayer and thus the attenuationof the thermally induced strains. It is also interesting to note that the residualstrains induced in the non-graded sample were most profound along theplanes parallel with the sharp interfaces, resulting in the concomitant crackingand sample disintegration. In contrast, the graded sample showed little or nocracking. This observation verifies the importance of designing bilayer ceramicswith graded interfaces to reduce the formation of undesirable residual strainsor stresses which may cause cracking and delamination at the interface.

As would be expected, the microhardness (Hv) of the graded materialincreases with depth from 12.1 to 14.3 GPa (Fig. 5.9), since the concentrationof relatively soft AT decreases with depth [Pratapa et al., 1998; Skala, 2000;Manurung, 2001]. The stress–strain curves obtained from spherical indentationson both graded and control samples are illustrated in Fig. 5.10. Solid curvesare empirical fits for the data. The data deviate from the Hertzian elasticlimit at stresses above Po ª 2 GPa for the control and Po ª 0.2 GPa for thegraded sample, marking the onset of ‘yield’. This result serves to verify the‘quasi-plastic’ nature of the graded layer, a phenomenon which has also beenobserved in other ceramics with a heterogeneous microstructure [An et al.,1996; Padture et al., 1995; Liu et al., 1996].

Figure 5.11 is an optical micrograph showing the Vickers indentationdamage around the indent at 200 N load. There is a distinct upheaval in thevicinity of the indent as a result of pronounced grain uplift. The presence ofa profuse damage zone surrounding the indent is vividly highlighted underthe Nomarski illumination. However, no radial cracks are observed at 200 N.Unlike the brittle control sample which exhibits cracks emanating from all

0 0.5 1 1.5 2 2.5 3Depth (mm)

25

20

15

10

5

0

Har

dn

ess

(GP

a)

AT rich AT poor

surface end

5.9 Vickers hardness of alumina–AT LGM as a function of depth withload of 3 kg. Error bars indicate two mean deviations (±).


four indentation corners, the graded surface of LGM exhibits either no cracks(up to 200 N) or short cracks in only one or two corners of the indent, at 300N. Grain pushout is routinely observed around the indent. It is believed thatduring the process of indentation, the weakly bonded grains are initiallydebonded, lifted up and eventually pushed out from their original positions.It appears that most of the indentation energy is used for debonding, liftingand pushing out the grains from the surface, thus rendering the materialdamage resistant. The high propensity for grain debonding is believed toarise from the presence of residual tensile stresses produced by the thermalexpansion mismatch between AT and alumina.

It should be pointed out that the display of profuse grain uplift andmicrodamage around the indent is unusual for ceramic materials. When anindentation is applied, ceramics usually exhibit a ‘sink in impression’ whichis caused by the densification below the tip of the indent [Zeng et al., 1996].This is usually accompanied by formation of radial cracks at the four corner-tips of the indents, indicating the brittleness of the material. By contrast,metals usually exhibit a ‘rising of material’ or surface uplift above theunindented surface level as a result of shear or plastic deformation. It followsthat the graded material exhibits indentation damage patterns that are indicativeof pseudo-plastic deformation during loading. The presence of AT ‘softens’the alumina matrix, thus rendering it effective in energy-absorption andcrack attenuation [Pratapa & Low, 1998; Skala, 2000].

Optical microscopy confirms that the ‘yielding’ behaviour during Hertzianloading arises from the onset of indentation damage in the graded region ofthe LGM. The nature of this damage can be discerned from the micrographsin Fig. 5.12, obtained using the bonded-interface section technique previouslydescribed. The micrographs show both half-surface (upper) and section views

0.00 0.03 0.06 0.09 0.12 0.15Strain

8

7

6

5

4

3

2

1

0

Str

ess

(GP

a)

5.10 Indentation stress–strain curves for both graded (�) and control(D) samples.


(lower) of indentations. The results clearly show the evolution of subsurfacedamage development as one progresses up the stress–strain curve in Fig.5.10 that corresponds to increasing indentation pressure. The initiation of thedeformation subsurface damage zone and subsequent expansion of this zoneare immediately apparent from the grain deformations or displacementsrevealed by the Nomarski contrast. At Po ª 0.5 GPa, i.e. just above the elasticlimit, only a few grains have deformed through shear-driven debonding,sliding and pushout. At increasing pressures, the number of deformed grainsincreases, and the damage zone expands further below the surface. At Po ª1.5 GPa, the damage is more profuse and begins to take on the appearanceof the well-developed, near-hemispherical deformation zone expected from

(a)

(b)

5.11 Indentation damage on the surface of LGM at a load of 200 N asrevealed by (a) Nomarski contrast and (b) SEM.

20 mm


continuum plasticity models [Johnson, 1985]. The complete absence of anyring cracks or cone cracks on the surfaces even at the maximum pressuresuggests that the graded material is damage tolerant. Cone fracture is inhibitedby the ability of the material to contain the extent of microdamage to a smallarea around the indent via multiple energy-absorbing mechanisms whichinclude diffuse microcracking, grain debonding and sliding, crack deflection,grain pushout, and grain bridging. In contrast, a ring crack on the surface ofthe control sample has initiated and attempted to run around the contactcircle. The classical cone crack has also formed in the subsurface region,highlighting the brittleness of this material.

LGMs of the AT/alumina and AT/ZTA displayed some very interestingproperties which include excellent machinability, low thermal expansioncoefficient, improved thermal shock resistance, low hardness (about 5 GPa),low Young’s modulus (E) (250 GPa) and excellent flaw tolerance [Pratapa,1997; Pratapa & Low, 1998; Skala, 2000; Manurung, 2001]. These materialsappeared to display a large degree of near-surface ‘quasi-plasticity’ underthe Hertzian or the Vickers indenter which effectively inhibits the formationand propagation of cracks. The ‘ductile’ behaviour of these materials was

(a) (b)

5.12 Optical micrographs in Nomarski illumination showing half-surface (top) and section (bottom) views of Hertzian damage in bothgraded and control samples: (a) control sample at P = 1.5 kN;(b) graded sample at P = 1.5 kN.


further verified by the dependence of Vickers microhardness (H) on indentationload and a low ratio of H to E. These materials also exhibited a superiorthermal shock resistance when compared to pure alumina.

No cracks were observed on the surface of these materials during theVickers or Hertzian indentation [Low et al., 1996b; Pratapa, 1997; Skala,2000]. Instead, a heavily microdamaged region was observed both withinand in the vicinity of the indenter. The formation of these damaged zones isbelieved to act as an effective ‘energy sink’ for the indenter, thus shieldingthe material from crack formation. Energy dissipative processes such asdebonding between AT and alumina grains, sliding and pushout of AT grainshave been observed in these damage zones. This display of flaw tolerance isattributed to high thermal expansion anisotropy of AT grains which inducevery large residual stresses. These stresses can cause alumina grains to form‘crack bridges’ and thus apply closure forces for shielding the crack tip fromthe applied stress intensity field, not unlike microcrack toughening [Runyan& Bennison, 1991]. Very similar observations in Vickers and Hertzian contactdamages were obtained for the layered-graded ZTA/AT system [Pratapa,1997; Pratapa & Low, 1998].

5.4.3 Alumina/mullite/AT hybrid

The Al2O3/mullite/AT hybrid has been fabricated by infiltrating an Al2O3

preform with both TiCl4 and Si(OC2H5)4, followed by heat-treatment at1600∞C for 2 h. The phase composition on the surface of the hybrid sampleas revealed by XRD showed the presence of alumina, mullite and AT. Nopeaks associated with that of rutile or silica were visible, which suggests thatcomplete reactions between titania and alumina were achieved to form AT(Al2TiO5) at ~1300∞C, and between silica and alumina to form mullite (3Al2O3

· 2SiO2) at ~1100∞C as follows:

TiO2 + Al2O3 Æ Al2TiO5 (5.7)

SiO2 + Al2O3 Æ 3Al2O3 · 2SiO2 (5.8)

Optical microscopy of the cross-section of an infiltrated hybrid sample revealeda distinct gradation of AT and mullite content in the infiltrated zone at themicrometre scale. However, whether this gradation occurs at the nanometrescale cannot be discerned from the microstructure. It is interesting to notethat the interface between graded and non-graded regions shows exaggeratedgrowth of alumina grains, which is not observed in the graded hybrid region.Scanning electron microscopy of the hybrid region showed the presence ofneedle-like mullite and equiaxed AT grains embedded within the finemicrostructure. A scanning electron micrograph of an as-fired surface of thehybrid sample showed the presence of needle-like mullite grains with adense interlocked structure. These elongated grains help to achieve flaw


tolerance in the hybrid through crack bridging and crack deflection as energydissipative processes. Small equiaxed or acicular AT grains are observed todistribute along or at the triple-point junctions of larger alumina grains. Thepresence of both mullite and AT has also resulted in much finer Al2O3 grainswithin the hybrid region, indicating their effectiveness as grain-growthinhibitors. Such grain refinement of Al2O3 has also been observed in Al2O3–mullite, Al2O3–AT and Al2O3–CaAl12O19 systems [Marple & Green, 1993;Skala, 2000; Asmi et al, 1999] and may account for the relatively highhardness observed. The combined effect of mullite and AT has allowed amuch smaller reduction in hardness when compared to the much softer Al2O3–AT system. Clearly, the self-reinforcement due to the presence of mullite hascompensated for the much softer AT phase.

XRD and grazing-incidence synchrotron radiation diffraction (GISRD)plots of a hybrid sample at different depths from the surface showed theabundance of a-Al2O3, AT and mullite to vary with depth (Fig. 5.13). As thedepth increased, the abundance of both mullite and AT decreased, but that ofaAl2O3 increased. The composition depth profiles as determined from theRietveld analysis are shown in Fig. 5.13(a). From the results it can be seen

MulliteATAl2O3

0 0.5 1 1.5 2 2.5Depth (mm)

(a)

100908070605040302010

0

Per

cen

tag

e o

f p

has

es

10 100 1000 10000Penetration depth (Å)

AT

M

1.6

1.4

1.2

1

0.8

0.6

0.4

0.2

0

Co

un

ts r

atio

(A

T/A

, M

/A)

5.13 Depth profiling of phase compositions in a hybrid LGM asrevealed by (a) X-ray diffraction and (b) grazing-incidencesynchrotron radiation diffraction.


that the abundances of AT and mullite in the hybrid sample decrease rapidlywithin the first 0.5 mm. The existence of graded compositions at the nano-and micro-scale has been established by GISRD. Figure 5.13(b) shows theGISRD plots of peak intensity ratios for AT and mullite relative to aluminaas a function of penetration depth (l) [Singh & Low, 2002a,b; Low et al.,2002. As the grazing-angle (a) or depth of penetration (l) increases, the ATand mullite peaks become less intense as compared to alumina. The ATintensity ratio curve shows a rapid decrease within the depths of <50 Å,above which there is a gradual fall. It is interesting to note that the mulliteintensity ratio curve shows a more gradual fall throughout the depths. Thisis indicative of a nanometre-scale gradation in phase composition within thehybrid material as would be expected from the time-dependent kinetics ofthe infiltration process. In essence, the near-surface hybrid layer is richer inboth AT and mullite, which decrease in abundance rapidly with depth.

Figure 5.14 shows the depth-profile variation of hardness (Hv) and fracturetoughness (KIC) of the hybrid sample. Fracture toughness is evidently highernear the surface due to the abundant presence of both AT and mullite. As thedepth increases, the abundance of both phases and thus the fracture toughnessdecreases. Hardness is lower near the surface due to the abundant presenceof much softer AT, but increases with depth towards the alumina region. Itshould be noted that since measurements were made on a polished cross-sectioned sample, the indentation data could not be collected at distancesless than 0.5 mm from the surface. Hence, hardness and fracture toughnessvalues on the surface were measured separately.

When compared to the control, the presence of AT and mullite as dispersedphases has imparted a two-fold increase in the fracture toughness withoutcausing a large reduction in hardness as in the alumina-AT system (Table5.3) [Singh & Low, 2002a; Low, 1998a, 1998b]. This implies the display ofboth self-strengthening and self-toughening processes in the hybrid sampleby virtue of the presence of both low thermal expansion AT and needle-like

0 0.5 1 1.5 2 2.5Distance from the surface (mm)

Hardness

Toughness

10

8

6

4

2

0

Frac

ture

to

ug

hn

ess

(MP

a.m

1/2 )

20

18

16

14

12

10

Hard

ness (G

Pa)

5.14 Variation of hardness and fracture toughness as a function ofdistance for the hybrid LGM.


mullite and its interlocked texture within the microstructure. The fracturetoughness of over 8.0 MPa.m1/2 is similar to that reported for the alumina–AT system but considerably higher than that of the alumina–mullite system[Padture et al., 1993; Chan et al., 1996; Low, 1998a, 1998b; Marple &Green, 1991]. The presence of elongated mullite grains acting as bridgingsites augmented by large residual stresses due to thermal expansion mismatchbetween alumina and AT is believed to be responsible for the much improvedfracture toughness. The characteristic elongated grain morphology of mulliteand an interlocked texture would have also resulted in desirable crack-tipdeflections and tilts to produce tortuous crack paths.

Figure 5.15 shows the KIC versus load plot for the graded region of thehybrid sample. A rising R-curve is clearly evident which suggests the presenceof flaw tolerance behaviour in the graded hybrid region. This display of flawtolerance may be attributable to crack-tip energy dissipative processes suchas crack bridging, crack deflection and crack branching by virtue of thesimultaneous presence of elongated mullite and soft but low thermal expansionAT grains.

5.5 Concluding remarks

The concept and synthesis of layered and graded materials (LGMs) from theinfiltration technique have been introduced. The mechanical and fracture

Table 5.3 Physical and mechanical properties of hybrid and control Al2O3 samples

Sample Shrinkage Apparent Bulk density Hv KIC

(%) porosity (%) (g.cm–3) (GPa) (MPa.m1/2)

Al2O3 19.5 0.25 3.93 17.9 (0.4)* 3.5 (0.4)Hybrid 6.7 1.08 3.63 16.9 (0.2) 8.3 (0.6)

*Figures in parentheses are the estimated standard deviation of the values to the left.

5.15 Display of crack-growth resistance in the hybrid LGM.

0 5 10 15 20 25 30 35Load (kg)

10

9

8

7

6

5

4

KlC

(M

Pa.

m1/

2 )


properties of various alumina-based LGMs have been discussed in relationto phase composition, microstructure and residual stress. These materialsdisplayed superior mechanical performance in terms of strength, fractureand thermal shock resistance. This new approach may be further developedfor designing new generation layer composites with an in-situ top wearresistant coating and a flaw-tolerant graded substrate or bulk. For instance,through unidirectional infiltration, a layer composite of a wear-resistant aluminasurface and a damage-tolerant graded AT/alumina bulk can be fabricated. Asimilar design can also be tailored for alumina– or ZTA–mullite systemswhere the in-situ top layer imparts wear and fatigue resistance.

The infiltration process also offers a potential for designing in-situ thinfilms of pristine HTSC with graded HTSC/epoxy substrates via careful controlof infiltrant kinetics from a specified direction. These structures will beideally suited for fabrication of superconducting electronic devices such asmicrobridges, microwave cavities and resonators. A similar avenue can beused to design graded ZrP/epoxy protonic conductors for use in a number ofelectrochemical devices such as fuel cells, gas sensors, and electrolysers. Anew philosophy and design process for tailoring layer graded composites forboth strength and damage tolerance has been highlighted for alumina-basedceramics. This approach allows the design of multifunctional new generationlayer composites with a top hard layer for wear resistance and a graded bulkor underlayer for damage tolerance. A similar design strategy can also betailored for other novel systems based on ceramic/polymer, ceramic/metal orpolymer/polymer hybridisation. This approach also has generic appeal andcan be extended to other more complex geometries or architectures (e.g.tubes or crucibles).

5.6 Acknowledgements

Financial support from the Australian Research Council (LX0242352 andA00001131) and the Australian Institute of Nuclear Science and Engineering(AINSE Grants 96/143 and 97/141, 98/010, 99/029 and 00/091P), the AustralianSynchrotron Research Program (ANBF 99/2000-AB-26 and 00/01-AB-31)and ISIS (RB12443 and RB13709) of the Rutherford-Appleton Laboratoryare acknowledged. IML is especially indebted to his past research students(R. Skala, D. Asmi, P. Manurung, S. Pratapa and D. Lawrence) and also theresearch colleagues (B. Lawn, D. Li, B. O’Connor, C. Buckley and M.Singh) for their invaluable contributions to the research program on layeredand graded materials over the last eight years.

5.7 References

An, L., Chan, H.M., Padture, N.P. & Lawn, B.R. (1996): Damage-resistant alumina-based layer composites. J. Mater. Res. 11, 204–210.


Asmi, D., Low, I.M., Kennedy, S.J. & Day, R.A. (1999): Characteristics of layered andgraded calcium hexaluminate/alumina composites. Mater. Lett. 40, 96–102.

Carman, P.C. (1956): Flow of Gaseous through Porous Media, Butterworths ScientificPublications, London.

Chan, H., An, L., Padture, N. & Lawn, B.R. (1996): Damage resistant alumina-basedlayer composites. J. Mater. Res. 11, 204–210.

Dullien, F.A.L. (1979): Porous Media Fluid Transport and Pore Structure. AcademicPress, New York.

Dullien, F.A., El-Sayed, M.S. & Batra, V.K. (1977): Rate of Capillary Rise in PorousMedia with Non-Uniform Porous, J. Colloids and Interface Science, 60, 497–506.

Einset, E.O. (1996): Capillary Infiltration Rate into Porous Media with Application ofSilicon Composite Processing. J. Am. Ceram. Soc., 79(2), 333–338.

Hirai, T. (1996): Chapter 20 in Processing of Ceramics, Part 2, edited. by R.J. Brook.VCH, Weinheim pp. 293–341.

Johnson, K.L. (1985): Contact Mechanics. Cambridge University Press, London.Koizumi, M. (1993): in Ceramic Transactions, Vol. 34 – FGMs (Proc. 2nd Int. Symp. on

FGMs), edited by J.B. Holt, M. Koizumi, T. Hirai and Z.A. Munir. Am. Ceram. Soc.,Westerville, OH, pp. 3–10.

Ligenza, J.R. & Bernstein, R.B. (1951): The Rate of Rice of Liquids in Fine VerticalCapillaries, J. Am. Chem. Soc. 73, 4636–4638.

Liu, H., Lawn, B.R. & Hsu, S.M. (1996): Hertzian contact response of tailored siliconnitride multilayers. J. Am. Ceram. Soc. 79, 1009–1014.

Low, I.M. (1998a): Processing of an in-situ layered and graded alumina-aluminium titanatecomposites. Mater. Res. Bull. 33, 1475–1482.

Low, I.M. (1998b): Synthesis and properties of in-situ layered and graded aluminiumtitanate/alumina composites. J. Aust. Ceram. Soc. 34, 250–255.

Low, I.M., Skala, R.D., Richards, R. & Perera, D.S. (1993): Synthesis and properties ofnovel mullite/ZTA composites. J. Mater. Sci. Lett. 12, 1985–1987.

Low, I.M., Skala, R. & Perera, D.S. (1994): Fracture properties of layered mullite/ZTAcomposites. J. Mater. Sci. Lett. 13, 1334–1336.

Low, I.M., Skala, R.D. & Zhou, D. (1996a): Synthesis of functionally-graded aluminiumtitanate/alumina composites. J. Mater. Sci. Lett. 15, 345–347.

Low, I.M., Skala, R.D. & Zhou, D. (1996b) Sol-gel processing of functionally-gradedceramics. Proc. Int. Workshop on Sol-Gel Processing of Advanced Ceramics, 8–9January 1996, Madras, India, Oxford & IBH Publisher, p. 143.

Low, I.M., Skala, R.D., Asmi, D., Manurung, P. & Singh, M. (2000): ‘Infiltration Processingof novel functionally-graded ceramic materials’. Pp 1464–1472 in Proc 2000 PowderMetallurgy World Congress (Eds. K. Kosuge & H. Nagai), 12–16 Nov, 2000, Kyoto,Japan

Low, I.M., Singh, M., Manurung, P., Wren, E., Sheppard, D.P. & Barsoum, M.W. (2002):Depth profiling of phase composition and texture in layered-graded Al2O3- and Ti3SiC2-based systems using x-ray and synchrotron radiation diffraction. Key EngineeringMaterials 224–226, 505–510.

Manurung, P. (2001): Microstructural design and characterisation of alumina/aluminiumtitanate composites. Ph.D Thesis, Curtin University of Technology, Perth, Australia.

Marple, B.R. & Green, D.J. (1990): J. Am. Ceram. Soc. 73, 3611.Marple, B.R. & Green, D.J. (1991): Mullite/alumina particulate composites by infiltration

processing: III, Mechanical properties. J. Am. Ceram. Soc. 74, 2453–2459.Marple, B.R. & Green, D.J. (1992): J. Am. Ceram. Soc. 75, 44.


Marple, B.R. & Green, D.J. (1993): Graded compositions and microstructures by infiltrationprocessing. J. Mater. Sci. 28, 4637–4643.

Padture, N.P., Bennison, S.J. & Chan, H.M. (1993): Flaw-tolerance and crack-resistanceproperties of alumina and aluminium titanate composites with tailored microstructures.J. Am. Ceram. Soc. 76, 2312–2320.

Padture, N.P., Pender, D.C., Wuttiphan, S. & Lawn, B.R. (1995): In-situ processing ofsilicon carbide layer structures. J. Am. Ceram. Soc. 78, 3160–3162.

Pratapa, S. (1997): Synthesis and character of functionally-graded aluminium titanate/ZTA ceramics. M.Sc Thesis, Curtin University of Technology, Perth, Australia.

Pratapa, S. & Low, I.M. (1998): Infiltration-processed functionally-graded AT/alumina–zirconia composites: II, Mechanical properties. J. Mater. Sci. 33, 3047–3053.

Pratapa, S., Low, I.M. & O’Connor, B.H. (1998): Infiltration-processed functionally-graded AT/alumina-zirconia composites: I, Microstructure and physical properties. J.Mater. Sci. 33, 3037–3046.

Runyan, J.L. & Bennison, S.J. (1991): Fabrication of flaw-tolerant aluminium-titanate-reinforced alumina. J. Eur. Ceram. Soc. 7, 93–99.

Sakai, M. & Hirai, T. (1991): Fabrication and properties of FGMs. J. Ceram. Soc. Japan.99, 1002.

Semlak, K.A. & Rhines, F.N. (1958): Trans. Met. Soc. AIME, 212, 325.Singh, M. & Low, I.M. (2002a): Depth-profiling of phase composition and preferred

orientation in a graded alumina/mullite/aluminium-titanate hybrid using x-ray andsynchrotron radiation diffraction. Mater. Res. Bull. 37, 1279–1291.

Singh, M. & Low, I.M. (2002b): Layered and graded alumina-based hybrid compositesby infiltration processing. Key Engineering Materials 224–226, 493–498.

Singh, M., Manurung, P. & Low, I.M. (2002): Depth profiling of near-surface informationin a functionally-graded Al2O3/Al2TiO5 composite using grazing-incidence synchrotronradiation diffraction. Mater. Lett. 55, 344–349.

Skala, R.D. (2000): Development of a functionally-graded alumina/aluminium-titanatecomposite. PhD Thesis, Curtin University of Technology, Perth, Australia.

Travitzky, N.A. & Shlayen, A. (1998): Microstructure and Mechanical Properties ofAlumina/Cu-O. Material Science and Engineering, A224, 154–160.

Washburn, E.W. (1921): Principles of Physical Chemistry, McGraw-Hill, New York.Yokota, M., Hara, A., Ohata, M. & Mitani, H. (1980): Trans. Jap. Inst. Metals, 21, 652.Zeng, K., Soederlund, E., Giannakopoulos, A.E. & Rowcliffe, D.J. (1996): Controlled

indentation: a general approach to determine mechanical properties of brittle materials.Acta Mater. 44, 1127–1132.

154

6.1 Introduction

In functionally graded materials (FGMs), the properties change graduallywith position/direction and are affected by composition, microstructure, porosityand so on. The usefulness of functionally graded composites with a gradedstructure concept was recognized by Shen and Bever in 1972.1 On the otherhand, systematic researches on production methods for functionally gradedmaterials were carried out around four years later by Niino and colleagues ofthe National Aerospace Laboratory at Sendai in Japan.2–4 However, increasedperformance of FGMs is hampered by increased manufacturing cost anddifficulties in production.5 Therefore, the usage of FGMs in the marketplacerequires an extensive search for cheaper, reproducible and reliablemanufacturing methods for FGMs.

6.2 Functionally graded materials

In the development of functionally graded materials, there are two approaches.One is to eliminate the boundary of laminated-type composites, therebyeliminating discontinuities in the properties at the boundary. The other optionis to make non-uniform distributions of dispersoids in a homogeneouscomposite, thus creating multiple functions within the material.2 Therefore,continuous variation in composition, microstructure and so on, results inchange in properties as a function of position in the component.6

FGMs offer increased efficiency with reduced weight and volume formotors and generators, giving a useful operating temperature to thermalbarrier materials by eliminating thermal stresses between the layers.Furthermore, FGMs have the potential for use in a wide range of engineeringapplications, for instance aerospace, high-efficiency engine components (e.g.piston heads, engine blocks), brake discs, hot gas valves and tubes, turbineengine components, fuel cells, certain piezoelectric devices and direct metaltools (e.g. injection molding tools and cutting tools) for industrial use,

6SiAlON based functionally graded materials

H M A N D A L, Anadolu University, Turkey andN Ç A L I S A C I K B A S, MDA Advanced Ceramics Ltd,

Turkey

SiAlON based functionally graded materials 155

biomaterials used in artificial human implants, drug delivery devices withrelease rate control, armor and armament components for defense and manymore.7,8 For these applications, control of material composition is necessaryboth for improving a variety of properties such as mechanical performance,toughness and strength, and for reducing interfacial stresses between dissimilarmaterials.

A variety of methods to produce graded materials have been reported inthe relevant literature. The most popular method is powder processingtechniques, where different mixtures of powders are stacked, slip cast, tapecast, infiltrated, electrochemically processed, spark plasma sintered (SPS)and so on.9,10 Fabrication of FGMs offers a technological challenge in large-scale graduation and requires elaborate processing facilities. Thus, increasedperformance of FGMs is mostly limited by increased manufacturing costs.Widespread transfer into the marketplace, therefore, requires an extensivesearch for cheap, reproducible and reliable manufacturing methods forFGMs.11,12

6.3 SiAlON ceramics

SiAlONs is a general name for a large family of the ceramic alloys based onsilicon nitride.13 They were first discovered independently at about the sametime (1971–72) in the United Kingdom at the University of Newcastle-upon-Tyne (by Jack and Wilson14) and in Japan (by Oyama15).

It is well known that silicon nitride is a covalently bonded material andhas a hexagonal structure. It can exist in two crystallographic modificationsdesignated a and b.16,17 Both structures are built up of corner-sharing SiN4

tetrahedra. The b-Si3N4 structure is obtained by ABAB… stacking of siliconand nitrogen atoms. This structure leads to continuous channels parallel tothe c-direction (Fig. 6.1). However, the a-Si3N4 structure is the result of anABCDABCD… stacking. The channels are thus closed and as a result, thereare two large interstitial sites in each unit cell, where cations can beaccommodated (Fig. 6.2). The unit cell parameters of the a phase are a =7.818 Å and c = 4.591 Å. The corresponding data for the b-modification area = 7.595 Å and c = 2.9023 Å, based on single crystal refinement.

There are two SiAlON phases that are most important for related engineeringceramics, a-SiAlON and b-SiAlON, which are solid solutions based on aand b-Si3N4 structural modifications, respectively. In single phase form, b-SiAlON has higher toughness (7–8 MPa m1/2), strength and thermalconductivity. On the other hand, a-SiAlON has excellent hardness (~20GPa) but slightly worse strength and toughness (3–4 MPa m1/2) compared tob-SiAlON ceramics. To combine the advantages of these phases, a/b-SiAlONcomposites have been developed.18 The other crystalline phase found inSiAlON systems is O-SiAlON, which has also drawn interest in the engineering


ceramics field. The ‘behaviour diagram’ of the SiAlON system at 1700–1730∞C is shown in Fig. 6.319 where it can be clearly seen that x-axis and y-axis scales are expressed in equivalent percentages of aluminum and oxygen,respectively. The five ordered SiAlON polytypoid phases (27R, 21R, 12H,15R and 8H) occur close to the AlN corner in this system and these phasesare structurally very similar to each other. The high-temperature reactions ofSi3N4 and Al2O3 usually give b-SiAlON and SiO2-rich X-phase.20 Additionsof suitable metal oxide (or nitride) will expand the SiAlON phase into a five-component system, Me-Si-Al-O-N, called a Jänecke prism.21

Silicon nitride is a highly covalent bonded compound with self-diffusioncoefficient of the nitrogen atoms of 6.3 ¥ 10–20 cm2/s at 1400∞C.22 Therefore,densification without any sintering additives is nearly impossible. In 1961,Deeley and Herbert.23 was the first to report that Si3N4 ceramics could be

x

y z

Silicon (IV) nitride

6.1 The b-Si3N4 structure is obtained by ABAB… stacking of siliconand nitrogen atoms.

Silicon

Nitrogen

6.2 The a-Si3N4 structure is the result of an ABCDABCD… stacking ofsilicon and nitrogen atoms.


densified by hot pressing with the incorporation of oxides as sintering additives.On the other hand, sintering additives limit the usage of Si3N4 at high-temperature applications. Thus, Oyama.15 and Jack and Wilson14 tried toovercome the problem of high-temperature strength degradation byinvestigating SiAlON ceramics.

The sintering behavior has a significant influence on the mechanical andthermal properties of sintered materials. Sintering of SiAlON ceramics iscarried out by a liquid-phase sintering process, which contains a residualgrain boundary phase inherited from the sintering medium. This phase canbe glassy (amorphous) or crystalline depending on factors such as overallcomposition and the applied cooling conditions. Liquid-phase sintering canbe performed in several ways, namely pressureless sintering, hot pressing(HP), hot isostatic pressing (HIP) and gas pressure sintering (GPS).

Pressureless sintering can be employed to fabricate complex shapes, andsignificant stabilization can be achieved by using a protective powder bedwhile the products generally indicate low density and the process requireslarge amounts of additives for densification.24

In hot pressing, powder mixtures of Si3N4 with additives are heated tohigh temperatures under an applied uniaxial pressure. Traditional hot pressinguses 20–30 MPa pressure which enhances both rearrangement of particlesand grain boundary diffusion. The hot pressing offers the ability to fabricatedense products, but also limits the products to simple shapes.25

Another method is to apply isostatic pressure and hot-isostatic pressing(HIP), now being another established technique. This technique is a veryattractive because it offers possibilities of making dense SiAlON ceramicswith a negligible residual glassy grain boundary phase and hence betterhigh-temperature properties. However, HP and HIP techniques are very costly.

6.3 The ‘behaviour diagram’ of the SiAlON system at 1700–1730∞C.

12H21H27H

Equivalent Al (%)

Si3N4 4(AIN)

2(Al2O3)

x-phase

15R

8H

b-sialon

a-sialon

3(SiO2)


They become economic only when a large number of samples is involved.26

Gas pressure sintering (GPS) is another technique to densify SiAlONceramics. GPS has proven to be an effective approach to minimize the contentof oxide additives at no sacrifice of density, because high nitrogen pressureis of advantage for suppressing decomposition of Si3N4 at temperaturesabove 1800∞C, and provides extra driving force for sintering.27,28

6.3.1 b-SiAlON ceramics

b-SiAlONs, which are the first developed group in the SiAlON materials,are formed by substituting up to two-thirds of the Si in the b-Si3N4 by Alprovided that valency compensation is maintained by the replacement of anequivalent concentration of N by O to give a range of b-SiAlONs,Si6–zAlzOzN8–z with 0 < z < 4.2.29 Thus, z (Si–N) bonds are replaced by z(Al–O) bonds, since the difference between the respective bond lengths(1.74 Å for Si–N and 1.75 Å for Al–O) is small and the extent of replacementis wide.

With increasing z value the density of b-SiAlON decreases linearly, similarlylowering Young’s modulus, strength, thermal conductivity, hardness andfracture toughness.30 On the other hand, for low z values (z < 1) the hardnessand fracture toughness increase and when z ≥ 1 they decrease. It has beenalso depicted that glass-free microstructures could be obtained in b-SiAlONpolycrystals with substitutional level z ≥ 2.31

Monolithic b-SiAlON materials can be easily densified by pressurelesssintering, since the presence of alumina in the starting mixture lowers theeutectic temperature of the densifying liquid phase by some 200–300∞C. Insingle-phase form, b-SiAlON materials possess a high degree of toughness(8 MPa m1/2) by in-situ reinforcement with the elongated b-SiAlON grains.They also possess good strength up to 1000 MPa and excellent thermalshock resistance.32

6.3.2 a-SiAlON ceramics

a-SiAlONs were first observed very shortly after b-SiAlONs with the generalformula of MxSi12–m–nAlm+nOnN16–n, where M is one of the cations Li, Mg,Ca or Y or most rare earths (excluding La, Ce, Pr and Eu); x is equal to mdivided by the valency of the M cation; there is a minimum value for the xparameter (0.3–0.5; i.e. there is an immiscibility gap between the a-Si3N4

and a-SiAlON phases); and x cannot exceed 2 since there are only twointerstitial sites in each unit cell. In a-SiAlONs, n (Si–N) bonds (1.74 Å) arereplaced by similar-sized Al–O (1.75 Å). The larger lattice strain resultingfrom the replacement of Si–N by Al–N restricts the range of solid solution(m-value) in a-SiAlON compared with b-SiAlON. The n-values are expectedto favor b-SiAlON formation.33


The a-SiAlON has twice the unit-cell parameter in the c-direction comparedto b-SiAlON, and the doubling of the Burgers vector for c[0001] dislocationsmeans that dislocation movement is more difficult and the hardness is enhanced.

Studies by Ekström and co-workers34 on the sintering of a-SiAlON ceramicsshowed that the largest rare earth cation to enter the a-SiAlON structurealone is Nd3+, with an ionic radius of 0.99 Å. The same authors also claimedthat the slightly larger cation Ce3+, with a radius around 1.03 Å, could notenter the a-SiAlON structure alone, but may enter when mixed with a smallerstabilizing cation like Y3+ with a radius of 0.89 Å.

a-SiAlONs were known to have a microstructure of fine and equiaxedgrains. But, more recently, it has been proved that a-SiAlON ceramics canalso be produced with elongated morphologies. The first example for theelongated nature of a-SiAlON grains was observed by Hwang et al.35 usinga mixture of CaO–SrO–Y2O3 as sintering additives. Therefore, fracturetoughness and hardness increase at the same time in the material.36

6.3.3 O-SiAlON ceramics

O-SiAlON is another crystalline phase of interest. There is a limited solubilityof alumina in silicon oxynitride structure to give O-SiAlONs, represented bythe formula Si2–xAlxO1+xN2–x, where x varies from zero to ~0.2. Formation ofO-SiAlON occurs in the same mechanism as b-SiAlON; i.e. Si + N is replacedby Al + O. The lattice parameters of O-SiAlON, Si2–xAlxO1+xN2–x, increasein a very typical way with the x value.37

The formation of Si2N2O from mixtures of high-purity Si3N4 and SiO2

heated at high temperatures is very sluggish due to kinetic hindrance. ForSi2N2O materials intended for high temperatures, addition of metal oxidessuch as Y2O3 alone, which forms a refractory glass, is the most attractivealternative.

The mechanical properties of monophase Si2N2O are difficult to determinebecause of preparation difficulties in obtaining a monophasic Si2N2O materialwithout the use of a sintering aid. A good approximation of Vickers hardness(HV 10) of 1600 kg/mm2 and a fracture toughness of ~3.3 MPa m1/2 wereobtained from measurements on the materials HIPed at 1900∞C, consistingof almost pure, crystalline Si2N2O.38 A slight improvement of the toughnesshas been found in O-SiAlON (x = 0.1) material with Y2O3 addition. Themodulus of rupture of similar O-SiAlON materials measured by Trigg andJack is ~420 MPa.39 Besides, O-SiAlON materials have poor toughnessdue to strong bonding Si2N2O crystals and the SiO2-rich glassy phase.O-SiAlON is characterized by high oxidation resistance due to its high oxygencontent.


6.3.4 a-b-SiAlON ceramics

a- and b-SiAlON phases are completely compatible and a-SiAlON/b-SiAlONcomposites are readily prepared by a single-stage sintering of the appropriatemixture of nitrides and oxides.40 Therefore, in recent years, mixed a-b-SiAlON materials have received increasing attention owing to their easierfabrication compared to Si3N4 materials. More importantly, good mechanicalproperties can be obtained due to the high hardness of a-SiAlON and thegood strength and toughness of b-SiAlON.41 However, the hardness andtoughness of a-b-SiAlON composites are not as high as those of theirmonolithic counterparts. Moreover, it has recently been found that the phasecomposition and microstructure of a- and a-b-SiAlON ceramics are greatlyaffected by post-sintering heat treatment procedures at lower temperatures(1300–1600∞C) when rare earth oxides are used as sintering additives. Thea-SiAlON phase is only stable at high temperatures and transforms to b-SiAlON plus other crystalline or vitreous phases.42–44 The ease with whichthis transformation proceeds decreases with increasing atomic number of therare earth cation. Indeed, in the absence of b-SiAlON nuclei, certain ytterbiuma-SiAlON compositions do not transform to b even when a high level of theliquid phase is present.45,46 This transformation provides a convenientmechanism for controlling the mechanical properties of the final material.However, it can only be used beneficially in applications where the maximumservice temperature is below the transformation temperature (1000∞C) of thegrain boundary glassy phase. High-temperature properties, especially oxidationand creep resistance, deteriorate significantly above this temperature.47

6.4 Functionally graded SiAlON ceramics

Microstructure, mechanical properties and high-temperature properties ofSiAlON ceramics can be optimized by a-b-SiAlON composites. Althougha-b-SiAlON composites possess better mechanical properties with respectto monolithic phases of either a- or b-SiAlON, a Æ b SiAlON phasetransformation restricts the compositional design of a-b-SiAlON ceramics.45

For this reason, the concept of producing functionally graded SiAlON ceramicshas been developed to improve especially the mechanical properties of thematerial. Depending on the application areas, the desired properties can beobtained with appropriately designed SiAlON composition and productionmethods. Development of functionally graded SiAlON ceramics providesbetter properties than either monolithic a-SiAlON or b-SiAlON, or a-b-SiAlON composites with continuous change in composition, microstructureand mechanical properties. Besides, a functionally graded structure is idealfor smooth reduction of thermal stresses in joints between the layers withdifferent composition.


Functionally graded SiAlON ceramics can be produced by lamination(powder and tape), powder bed, controlling the sintering conditions, infiltration,slip casting and so on. Some of these production techniques will be discussedin the following sections.

6.5 Production techniques of functionally graded

SiAlON ceramics

The aim of the preparation of a FGM is to achieve a well-controlled distributionof composition, microstructure and so on. Although many production techniquesfor functionally graded materials were already developed in the early 1990s,there are still limitations concerning material combinations, specimen geometryand cost. Thus, new processing techniques for producing functionally gradedmaterials are necessary.

6.5.1 Lamination technique

Lamination of green compacts or sheets, which are produced through thetape casting method, is a very important technique for graded materialproduction. For this purpose, layers with different a:b SiAlON ratio are usedto obtain graded structures. Further details regarding these productiontechniques are given below.

Powder lamination

Prelamination of green compacts is a simple and well-established techniquefor graded material production.48 In the powder lamination method, FGMfabrication stages include compositional design, compaction of differentpowders, and sintering at suitable conditions.

In compositional design of functionally graded SiAlON ceramics, selectionof cations and a:b phase ratios are very important parameters. Type of cationsaffects the densification and diffusion between the layers. They also controlthe thermomechanical properties of SiAlON ceramics.

As mentioned before, a-SiAlON possesses high hardness, wear resistanceand oxidation resistance while b-SiAlON has high fracture toughness andgood flexural strength. Thus, the a:b ratio of a designed FGM is very importantaccording to its application area.

In the compaction stage, two or more different powder compositions arestacked by sequential filling of the die and uniaxially pressed in an automaticpress. Packing of powders and grain size are effective parameters in controllingthe composition and thus the properties. After compaction, the laminates aresintered to obtain a gradual structure.

Shen and Nygren reported on the production of functionally graded SiAlON


ceramics using the SPS method.49 However, no detailed information wasgiven in this reference or elsewhere. Besides, our research group has producedlaminar-type functionally graded SiAlON ceramics with a transition zonethickness of ~1.3 mm.50,51,52 In these studies,50 the effects of type of cations,amount of liquid phase, sintering temperature and time on the diffusionbetween the layers were investigated. For this purpose, two-layered functionallygraded SiAlON ceramics were produced. One layer consisted of thecomposition with about equal proportions of a and b (B1, B2 and B3 rich inthe b-phase) and another layer was rich in the a-phase (A1 composition).The difference between B1 and B2 is in the type of cations: B1 contains onlya single cation while B2 has three types. The difference between B2 and B3is in the amount of the liquid phase: B3 contains more liquid phase.

Laminates were prepared by sequential filling of the die and pressing, asschematically shown in Fig. 6.4. The laminates B1–A1 (named FGM-A),B2–A1 (named FGM-B) and B3–A1 (named FGM-C) were sintered under22 bar nitrogen gas pressure at 1800∞C for 1 hour. Also, to observe theeffects of sintering time and temperature, B3–A1 laminate was sintered at 22bar nitrogen gas pressure at 1700∞C for 1 hour (named FGM-D) and for 2hours (named FGM-E).

The phase composition from the surface to the interior of the samples wasdetermined by X-ray diffractometry (XRD) through successive grinding ofthe surface at 100 mm intervals. The microstructural characterization of thesintered specimens was achieved by scanning electron microscopy (SEM) inbackscattered mode. The hardness change from the surface to the interior ofthe sample was measured by the Vickers indentation method at 19.6 N load.

Visual examination showed that laminates other than B1–A1 wereintermixed, showing that counter-diffusion occurred. SEM analyses alsoconfirmed the visual examinations of B1–A1 laminate (Fig. 6.5). It is obviousthat a sharp transition zone exists. This can be explained by the fact that thetype of cations affects the diffusion between the layers.

For B2–A1 laminate, there is a gradual transition from an a-rich region tob-rich one in XRD analyses. The transition zone is about 400 mm in thickness,and the b-SiAlON proportion in this zone changed from 20% to 70% asexpected (Fig. 6.6). The gradual change in the amount of phases was also

6.4 Basic demonstration of powder lamination method.

A1

B1


confirmed by SEM (Fig. 6.7) in that counter-diffusion of species had takenplace between B2 and A1. XRD measurements and microstructural observationscoincided with the hardness measurements, as shown in Fig. 6.8. Hardnessof the intermediate zone increased from 15 GPa to 19 GPa in moving fromthe b-SiAlON-rich side to the a-SiAlON-rich side.

Another combination was prepared by using powders of B3 and A1. Thislaminate also showed a gradual transition zone similar to the B2–A1 pair

6.5 BE-SEM analyses of B1–A1 laminate.

0 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 1.8 2Distance from surface (mm)

90

80

70

60

50

40

3020

10

0

% a

-SiA

lON

A1

B2

Transitionzone

6.6 Change in a-SiAlON ratio of sintered B2–A1 laminate ground upto 100 mm through the sample thickness.

50 mm 50 mm 50 mm

B2 A1Transitional zone

6.7 BE-SEM analyses of B2–A1 laminate.


(Fig. 6.9). Hardness values increased gradually from 16 GPa to about19 GPa, shifting from the b-SiAlON-rich region to the a-SiAlON-rich one(Fig. 6.10). However, the gradient layer was longer in this sample, whichmay be due to easier diffusion of species in a larger amount of liquid phasein B3.

In addition, the effect of sintering time and temperature on the diffusionbetween the layers was investigated for B3–A1 laminate. Changes in theamount of a-SiAlON from the surface of the samples sintered at differenttemperatures for different times (designated FGM-C, D, E) are shown in Fig.6.11. A gradual transition from a- to b-rich regions for each FGM wasobserved. Sintering time and temperature did not appear to affect the diffusionbetween the layers. Thus, sintering at 1700∞C for one hour (FGM-D) wassufficient. The gradual change in the amount of the phases was also confirmedby the SEM observations for each functionally graded SiAlON ceramic. Themicrostructure of each of the FGMs produced looked similar. A representative

A1

B2

Transition zone


20

19

18

17

16

15

14

Har

dn

ess

(GP

a)

6.8 Change in the hardness measurements of B2–A1 laminates alongthe sample thickness.


A1

B3

Transitionzone

90

80

70

60

50

40

30

20

10

0

% a

-SiA

lON

6.9 Change in a-SiAlON ratio of sintered B3–A1 laminate ground upto 100 mm through the sample thickness.


SEM image is given in Fig. 6.12. Furthermore, adding the same dopants inboth layers eliminated the sharp transition zone. Along the transition zone,needle-like a-SiAlON grains were observed. This confirms that the amountof glassy grain boundary phase is high enough for needle-like a-SiAlONformation and easy diffusion between the layers.

Both XRD and SEM studies support the results of hardness measurement,as illustrated in Fig. 6.13. Hardness values decreased gradually from ~19 to~15 GPa, over a distance from a-SiAlON to b-SiAlON rich sides for eachfunctionally graded SiAlON ceramic.

The results obtained revealed that the gradient layer was longer for theB3–A1 laminate than for the B2–A1 laminate.50 This can be explained bythe fact that the larger amount of liquid phase in the B3 layer facilitates thediffusion between the layers. The resulting functionally graded SiAlONceramics may be potential candidates for wear applications.

B3

A1

Transitionzone

FGMCFGMDFGME


90

80

70

60

50

40

30

20

10

0

%a-

SiA

lON

6.10 Change in the hardness measurements of B3-A1 laminatesalong the sample thickness.

A1

B3

Transitionzone


19

18.5

18

17.5

17

16.5

16

15.5

Har

dn

ess

(GP

a)

6.11 Change in the amount of a-SiAlON from the surface of thesamples sintered at different temperatures for different times (namedFGM-C, D, E).


Tapes lamination

Tape casting is a process for producing thin single-layer or stacked andlaminated multilayer structures for which adequate thickness control isrequired.53 Although there are some publications on the production of Si3N4

ceramics by the tape casting method,54,55 there is only one report availableon the preparation of SiAlON sheets by tape casting.56 According to thisstudy, Xu and co-workers prepared a-SiAlON sheets by tape casting andpressureless sintering at 1750∞C for 2 hours.

Preparation of ceramic bodies by tape casting has been considered to be

~ 50 mm

~ 50 mm ~ 50 mm ~ 50 mm

~ 50 mm~ 50 mm80a:20b

20a:80b(a) (b) (c)

(d) (e) (f)

6.12 BE-SEM analyses of FGM-C laminate.

FGM CFGM DFGM E

B3

A1


20

19

18

17

16

15

14

Har

dn

ess

(GP

a)

6.13 Change in the hardness measurements of FGM-C, D, Elaminates along the sample thickness.


one of the most favorable and visible routes in functionally graded materialsproduction.57–61 According to these studies, tape casting is a candidate toproduce multifunctional materials due to an adjustable layer thickness ofeach component. When the variations in compositions between layers arerelatively small, the material approaches a continuous composition, and stressconcentrations at the interfaces can be decreased to practically insignificantvalues.

In the literature, there has been no report about production of functionallygraded SiAlON ceramics by tape casting. The main advantage of this methodwith respect to others is that continuous change in composition, microstructureand mechanical properties can be obtained by stacking controlled layerthicknesses of different tape compositions.

For the production of functionally graded SiAlON ceramics, five differentSiAlON compositions were designated for tape casting. These are 85a:15b,70a:30b, 55a:45b, 40a:60b and 25a:75b from surface to bottom.62 Thetape casting was carried out with a laboratory tape caster on a glass substrate,as schematically indicated in Fig. 6.14. The blade gap was adjusted to 400mm. After drying at room temperature, sheets were cut 1 ¥ 1 cm in dimensions.Lamination was performed by compression of the green tapes with differentcompositions, where the surface was rich in a-SiAlON and the bottom richin b-SiAlON (5 sheets ¥ 5 compositions = 25 layers) under ~7 MPa pressureat room temperature (Fig. 6.15). The laminates were further cold isostatically

Direction ofmovement

Wet tape

Slurry

Substrate

ABC

High a-SiAlON content

High b-SiAlON content

6.14 Schematic illustration of tape production.

6.15 Stacking of tapes with different compositions.


pressed under 300 MPa followed by a binder burnout stage.After drying, crack-free tapes could be easily removed from the substrate;

the green tapes exhibit smooth surfaces, good flexibility and adequate strength(Fig. 6.16). The thickness of the dried sheet was measured to be about 150mm. The tapes were sintered in a gas pressure furnace at 1800∞C for 60 minunder 22 bar nitrogen gas pressure.

Change in the amount of a-SiAlON for FGM-F from the top surface tothe bottom is shown in Fig. 6.17. This clearly illustrates that a-SiAlONcontent changes gradually through the thickness of the sample. Hardnessmeasurements showed high hardness at the surface (top) and low hardness atthe bottom (Fig. 6.18). By comparison of the hardness measurements andscanning electron microscopy–secondary electron mode (SE–SEM)observations (Fig. 6.19), gradual change in composition and good connectivitybetween the layers were observed. This indicates that compaction pressureof uniaxial pressing is high enough to provide connection between the layers,providing easy removal of binders during the burnout stage without disrupting

6.16 Photograph of green tape.

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 1.1 1.2Distance from surface (mm)

90

80

70

60

50

40

30

20

10

0

% a

-SiA

lON

6.17 Change in the amount of a-SiAlON from the surface of thesamples FGM-F.


the physical integrity of the compact. Microstructural observations verifythe effect of further cold isostatic pressing to achieve homogeneity in thesample. As can be seen from the representative SEM image of FGM-F (seeFig. 6.20), tape casting is a good candidate for obtaining continuous changesin composition, microstructure and thus mechanical properties through thematerial thickness. The internal stresses can be minimized relative tomultifunctional materials produced by other methods. This commonly leadsto improved material performance by the tape casting method. Thus, producingSiAlON ceramics through the tape casting method may extend the applicationareas of SiAlON ceramics.

0 0.2 0.4 0.6 0.8 1 1.2 1.4Distance from surface (mm)

Har

dn

ess

(GP

a)

18

17.5

17

16.5

16

15.5

15

14.5

200 mm

6.18 Hardness measurements of designed FGM-F through thethickness of the sample.

6.19 SE-SEM analyses of FGM-F.


6.5.2 Powder bed technique

Previously, Chen and co-workers developed graded in-situ SiAlON ceramicsby embedding b-SiAlON green compacts in an a-SiAlON powder bed.63

Their results showed that compositions, microstructures and properties ofthe graded SiAlON ceramic change gradually from the hard a-SiAlON withspherical morphology on the surface to the tough and strong b-SiAlON withelongated grains in the core, but the graded layer was limited to ~300 mm.Recently, Jiang and Kang developed a technique for in-situ formation of ana-SiAlON layer on a b-SiAlON surface.64 This technique consists of packinga compact of b-SiAlON composition in an a-SiAlON powder bed. Theyfound that it was possible to control the thickness of the a-SiAlON-richlayer by changing the presintering conditions during heating to sinteringtemperature. Thickness of the a-SiAlON layer increased with increase at thepresintering temperature. However, this can lead to grain growth and resultantdecrease in strength. Thus, alternative methods must be developed.

(a)

(b)

(c)

85a:15b

25a:75b

6.20 BE-SEM analyses of FGM-F.


In another study, Mandal and co-workers produced functionally gradedSiAlON ceramics using the powder bed method.50 In their study, b-SiAlONcompacts were embedded in two different homogeneously mixed powderbed compositions, a-SiAlON (100 wt%) and AlN:BN (50:50 wt%). Theeffects of powder bed composition and pressure on the formation of a-SiAlON on the compact surface were investigated.

For the powder bed method, a composition rich in b-SiAlON (named B2)was selected as a compact composition to observe compact–powder bedinteraction. Two different powder bed compositions, a-SiAlON (100 wt%)and AlN:BN mixture (50:50 wt%), were prepared. b-SiAlON-rich pelletswere embedded into the powder bed compositions, as schematically shownin Fig. 6.21. Both green and sintered pellets were embedded into the samepowder bed composition in order to compare the effect of presintering on theinteraction. Sintering of the pellets was carried out under 22 bar nitrogen gaspressure at 1800∞C for 1 hour. To understand the effect of pressure on theinteraction zone, pressureless sintering was also carried out for comparison.

Changes in the amount of b-SiAlON from the surface of the samples forgreen and presintered compacts are shown in Figs 6.22 and 6.23, respectively,after sintering in a-SiAlON and AlN–BN powder beds. Both figures clearlyillustrate that the AlN–BN powder bed is more effective for the formation ofa-SiAlON at the surface than the a-SiAlON powder bed. Formation of an a-SiAlON-rich layer on the component surface is due to the transfer of a-SiAlON-forming ions from the powder bed. This transfer could be in various

BN crucible

Pellet

Powder bed

6.21 Basic demonstration of powder bed method.

AlN:BNa-SiAlON

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1Distance from surface (mm)

908070605040302010

0

% b

-SiA

lON

6.22 Change in b-SiAlON ratio of green samples sintered in differentpowder beds.


ways: (i) solid-state reaction of the powder bed with the compact surface;(ii) formation of a liquid in the powder bed and its reaction with the surface;(iii) evaporation from the powder bed and condensation on the compact. Thefirst mechanism appears to be unlikely as the contact area of fine powder ofthe loose powder bed should be rather low for effective solid-state reaction.The second mechanism is not applicable to an AlN:BN powder bed as noliquid formation is expected. For a-SiAlON powder bed, liquid formationshould occur but not as much as one can expect due to the loose powder bed.Therefore, it is likely that the material transfer occurs via vaporization ofreactants from the powder bed and condensation on the surface of the compactfor reaction. It is possible that AlN evaporates from the AlN:BN powder bedwhich then reacts with the compact surface. Consequently, the surfacecomposition becomes richer in AlN, causing a-SiAlON formation. Sinceless AlN is available from the a-SiAlON powder bed, its effect on a-SiAlONformation at the compact surface is naturally less.

Figure 6.24 gives a comparison of the amount of b-SiAlON change of thegreen and presintered compacts in AlN:BN powder. The green compacts had

AlN:BNa-SiAlON


908070605040302010

0

% b

-SiA

lON

6.23 Change in b-SiAlON ratio of presintered samples sintered indifferent powder beds.

GreenPresintered


90807060

504030

2010

0

% b

-SiA

lON

6.24 Comparison of the amount of b-SiAlON on the surfaces of thegreen and presintered compacts in AlN-BN powder bed.


more interaction with the powder bed than the presintered compacts. This ispresumably due to easy diffusion of condensed phases through the less densestructure of b-SiAlON just before reaching full densification. The maximumthickness of the gradient zone is limited to 400 mm.

The effect of nitrogen gas pressure during sintering on the thickness of thediffusion zone is determined in the AlN:BN powder bed by using 1 bar and22 bar gas pressure. The results indicated that change in gas pressure duringsintering has an influence on the thickness of the gradient zone (Fig. 6.25).The smaller gradient zone thickness at 1 bar can be attributed to sweeping ofevaporating species away by flowing nitrogen gas used during sintering.

As a result, it is obvious that the powder bed technique is not as effectivefor the production of functionally graded SiAlON ceramics as the laminationmethod.

6.5.3 Controlling the sintering conditions

Previously, Mandal and co-workers obtained a gradual change of a-SiAlONcontent from the surface through the core by rapid cooling.42 However, theaim was different from the functionally graded material concept. Recently,they have used this technique in production of functionally graded SiAlONceramics. The achievement of gradual changes is explained as a function ofa Æ b SiAlON transformation.

In the fast cooling method, two different samples were heated and cooledvery rapidly. The first was B2 (same composition in powder laminationmethod); the second was B4, obtained from B2 by sintering in an AlN:BNpowder bed. The surface of B2 was ground until 70% b-SiAlON was obtainedbefore fast cooling treatment. Then, the specimens were inserted in a furnacewith high-speed cooling. They were heated to 1600∞C at a rate of 15∞C/minand cooled rapidly by immediate removal from the furnace.

Sintered under 22 barSintered under 1 bar


90

8070

605040

3020

100

% b

-SiA

lON

6.25 Comparison of the amount of b-SiAlON on the surfaces of greencompacts sintered under 22 bar and 1 bar in AlN:BN powder bed.


The B2 sample was treated to investigate the effect of the fast coolingmethod on the surface modification. There was no gradient layer after fastcooling on the surface of B2. This can be explained by the stability of b-SiAlON, not transforming to a-SiAlON on heating. When the B4 sample wasused in fast cooling experiments, an ~ 300 mm gradient zone was obtained atthe surface (Fig. 6.26). This is probably due to its modified surface compositionduring sintering in the AlN:BN powder bed becoming more transformable.

As a result, this method is not practical, since gradient zone thickness islimited to about 300 mm and the specimen may crack during rapid cooling.

It is evident that the gradient zone thickness on the surface varies between200 and 400 mm, but is limited to about 400 mm maximum in the powder bedand fast cooling methods. Considering that ceramic products usually need amachining operation to achieve near-net shape, and the amount of materialremoval during machining is about a few hundred microns depending on theapplication, neither method appears to be applicable. Furthermore, the methodsare not practical in that the powder bed needs to be renewed at each sinteringcycle and the specimen may crack during rapid cooling. Therefore, thelamination method gives more promise for functionally graded SiAlONs.This is because one can adjust physically the layer thickness as required, e.g.depending on machining allowance.


Since the macroscopic properties of a material are strongly dependent on itsstructure, studies of structure/property relationships are among the mostimportant issues in material science. The properties of SiAlON ceramics arestrongly influenced by their microstructure and chemical composition.

Although the functionally graded materials concept has been known since1972, the production of functionally graded SiAlON ceramics is still a subjectto explore. Especially, the formation mechanisms of the functional gradientstructure have not yet been fully understood. If diffusion mechanisms areexplored, compositional design would be easy and the thickness of the graded

6.26 Change in b-SiAlON ratio in B1 sample ground up to 70% b-SiAlON content as a function of fast cooling.

0 0.05 0.1 0.15 0.2 0.25 0.3 0.35 0.4Distance from surface (mm)

908070605040302010

0

% b

-SiA

lON


structure could be adjusted. In addition, by using different SiAlONcompositions, new materials with superior properties over either monolithicb-SiAlON or a-SiAlON and a-b-SiAlON composites can be developed forindustrial applications.

6.7 References

1. Shen, M., Bever, M.B. (1972), ‘Gradients in polymeric materials’, J. Mat. Sci., 7,741–746.

2. Hirai, T. (1996), Materials Science and Technology, A Comprehensive Treatment,Vol. 17B, ed. R.J. Brook, VCH Verlags, Weinheim, Germany.

3. Holt, J.B., Koizumi, M., Hirai, T., Munir, Z.A. (1993), Ceramic Transactions, Am.Ceram. Soc., Westerville, OH.

4. Damzik, R.J., Neubrand, A., Rödel, J. (2000), ‘Functionally graded materials byelectrochemical processing and infiltration: application to tungsten/copper composites’,J. Mat. Sci., 35, 477–486.

5. Suresh, S., Mortensen, A. (1997), ‘Functionally graded metals and metal ceramiccomposites: thermomechanical behaviour’, Int. Mater. Rev., 42(3), 85–116.

6. Koizumi, M. (1997), ‘FGM activities in Japan’, Composites, Part B, 28B, 1–4.7. Shin, H.K., Natu, H., Dutta, D., Mazumder, J. (2003), ‘A method for the design and

fabrication of heterogeneous objects’, Materials and Design, 24, 339–353.8. Ichikawa, K. (2001), Functionally Graded Materials in the 21st Century, Kluwer

Academic, Boston, MA.9. Marple, B.R., Boulanger, J. (1994), ‘Graded casting of materials with continuous

gradients’, J. Am. Ceram. Soc., 77, 2747–2750.10. Kieback, B., Neubrand, A., Riedel, H. (2003), ‘Processing techniques for functionally

graded materials’, Mater. Sci. Eng., A362, 81–105.11. Parameswaran, V., Shukla, A. (2000), ‘Processing and characterization of model

functionally graded materials’, J. Mat. Sci., 35, 21–29.12. Carrillo-Heian, E.M., Carpenter, R.D., Paulino, G., Gibeling, J.C., Munir, Z. (2001),

‘Dense layered molybdenum disilicide–silicon carbide functionally graded compositesformed by field-activated synthesis’, J. Am. Ceram. Soc., 84, 962–968.

13. Jack, K.H. (1976), ‘Review: sialons and related nitrogen ceramics’, J. Mater. Sci.,11, 1135.

14. Jack, K.H., Wilson, W.I. (1972), ‘Ceramics based on the Si–Al–O–N and relatedsystems’, Nature (London) Phys. Sci., 238, 28–29.

15. Oyama, Y. (1971), ‘Solid solution in the ternary system Si3N4–AlN–Al2O3’, Jpn. J.App. Phys., 10, 1687.

16. Grun, R. (1979), ‘The crystal structure of b-Si3N4’, Acta Crystallogr, B35, 800–804.17. Kohatsu, I., McCauley, J.W. (1974), ‘Re-examination of the crystal structure of a-

Si3N4’, Mater. Res. Bull., 9, 917–920.18. Ekström, T., Ingelström, I. (1986), ‘Characterization and properties of SiAlON

materials’, in Proc. Int. Conf. Non-oxide Technical and Engineering Ceramics, ed S.Hampshire. Elsevier Applied Science, London, 231–253.

19. Ekström, T., Nygren, M. (1992), ‘SiAlON ceramic’, J. Am. Ceram. Soc., 75, 259–276.

20. Thompson, D.P. (1989), Preparation and Properties of Silicon Nitride Based Materials,ed. D.A. Bonnell and T.Y. Tien, Materials Science Forum, Vol. 47, Trans TechPublications, Aedermannsdorf, Switzerland.


21. Jänecke, E.Z. (1907), ‘Über eine neue Darstellungsform der van’t HoffschenUntersuchung über ozeanische Salzablagerungen’, Z Anorg. Chem., 53, 319–326.

22. Kijama, K., Shirasaki, S. (1976), ‘Nitrogen self diffusion in silicon nitride’, J. Chem.Phys., 65, 2668.

23. Deeley, G., Herbert, J. (1961), ‘Dense Silicon Nitride’, Powder Metall., 8, 145.24. Munakata, H., Hayashi, T., Suzuki, H., Saito, H. (1986), ‘Presureless sintering of

Si3N4 with Y2O3 and Al2O3’, J. Mat. Sci., 21, 3501–3508.25. Hwang, S.L., Chen, I.W. (1994), ‘Reaction hot pressing of a and b-SiAlON ceramics’,

J. Am. Ceram. Soc., 77, 165–171.26. Olsson, P.O., Ekström, T. (1990), ‘HIP sintered b and mixed a-b SiAlONs densified

with Y2O3 and La2O3 additions’, J. Mat. Sci., 25, 1824–1832.27. Li, H.X., Sun, W.Y., Wang, P.L., Yan, D.S., Tien, T.Y. (1997), ‘The effect of GPS

parameters on mechanical properties of Y-a-SiAlON ceramics’, Ceramics International,23, 449–456.

28. Biswas, S.K., Riley, F.L. (2001), ‘Gas pressure sintering of silicon nitride – currentstatus’, Mat. Chem. Phys., 67, 175–179.

29. Kuwabara, M., Benn, M., Riley, F.L. (1980), ‘The reaction hot-pressing of compositionsin the system Al–Si–O–N corresponding to (b-SiAlON)’, J. Mat. Sci., 15, 1407–1416.

30. Haviar, M., Johannesen, O. (1988), ‘Unit cell dimensions of b-SiAlON’, Adv. Ceram.Mater., 3, 405–407.

31. Pezzotti, G., Kleebe, H., Okamoto, K., Ota, K. (2000), ‘Structure and viscosity ofgrain boundary in high purity SiAlON ceramics’, J. Am. Ceram. Soc., 83, 2549–2555.

32. Kishi, K., Umebayashi, S., Tani, E. (1990), ‘Influence of microstructure on thestrength and fracture toughness of b-SiAlON’, J. Mat. Sci., 25, 2780–2784.

33. Hampshire, S., Park, H.K., Thompson, D.P., Jack, K.H. (1978), ‘a-SiAlON ceramics’,Nature, 274, 880.

34. Ekström, T. (1993), ‘SiAlON ceramics sintered with yttria and rare earth oxides’, inMaterials Research Society Symposium Proceedings, 121.

35. Hwang, C.J., Susintzky, D.W., Beaman, D.W. (1995), ‘Preparation of multication a-SiAlON containing strontium’, J. Am. Ceram. Soc., 78, 588.

36. Mandal, H. (1999), ‘New developments in a-SiAlON ceramics’, J. Eur. Ceram.Soc., 19, 2349–2357.

37. Jack, K.H. (1973), ‘Nitrogen ceramics’, Trans J. Br. Ceram. Soc., 72, 376.38. Ekström, T., Holmström, M., Olsson, P.O. (1991), ‘Yttria doped Si2N2O’, in Proc.

4th Int. Symp. on Ceramic Materials and Components for Engines, Amsterdam.39. Trigg, M.B., Jack, K.H. (1988), ‘The fabrication of O-SiAlON ceramics by pressureless

sintering’, J. Mat. Sci., 23, 481–487.40. Ekström, T. (1989), ‘Effect of composition, phase content and microstructure on the

performance of yttrium SiAlON ceramics’, Mat. Sci. Eng., A109, 341–349.41. Cao, G.Z., Metselar, R., Ziegler, G. (1992), ‘Microstructure and properties of mixed

a-b SiAlONs’, in Proc. 4th Int. Symp. on Ceramic Materials and Components forEngines, Amsterdam, p. 188.

42. Mandal, H., Thompson, D.P., Ekström, T. (1993), ‘Reversible a´b phasetransformation in heat treated SiAlON ceramics’, J. Eur. Ceram. Soc, 12, 421.

43. Ekström, T., Shen, Z.J. (1995), ‘Temperature stability of rare earth doped a-SiAlONceramics’, in Proc. 5th Int. Symp. on Ceramic Materials and Components for Engines,p. 206.


44. Zhao, R., Cheng, Y.B. (1996), ‘Decomposition of Sm a-SiAlON phases during postsintering heat treatment’, J. Eur. Ceram. Soc., 16, 1001.

45. Camuscu, N., Thompson, D.P., Mandal, H. (1997), ‘Effect of starting composition,type of rare earth sintering additive and amount of liquid phase on a ´ b phasetransformation’, J. Eur. Ceram. Soc., 17, 599.

46. Mandal, H., Thompson, D.P., Liu, Q., Gao, L. (1997), ‘High temperature stability ofa-SiAlON ceramics containing glass additions’, Eur. J. Solid. State. Inorg. Chem.,34, 179.

47. Klemm, H., Hermann, M., Reich, T., Schubert, C. (1998), ‘High temperature propertiesof mixed a-b SiAlON materials’, J. Am. Ceram. Soc., 81, 1141–1148.

48. Rabin, B.H., Heaps, R.J. (1993), in Holt, J.B., Koizumi, M., Hirai, T., Munir, Z.A.(eds), Ceramic Transactions, Vol. 34, Functionally Gradient Materials, Am. Ceram.Soc., Westerville, OH, pp. 173–180.

49. Shen, Z., Nygren, M. (2002), ‘Laminated and functionally graded materials preparedby spark plasma sintering’, Key Eng. Mater., 206–213, 2155–2158.

50. Çalıfl, N., Kuflhan, fi.R., Kara, F., Mandal, H. (2004), ‘Functionally graded SiAlONceramics’, J. Eur. Ceram. Soc., 24, 3387–3393.

51. Çalıfl, N., Kara, A., Kara, F., Mandal, H. (2004), ‘Development of laminar typefunctionally graded SiAlON ceramics’, Key Eng. Mater., 264–268, 1095–1098.

52. Çalıfl, N. (2002), ‘Functionally graded SiAlON ceramics’, BSc Thesis, AnadoluUniversity, Turkey.

53. Mistler, R.E., Twiname, E.R. (2000), ‘Tape casting theory and practice’, Am. Ceram.Soc., Westerville, OH, pp. 62–82.

54. Bitterlich, B., Heinrich, J.G. (2002), ‘Aqueous tape casting of silicon nitride’, J Eur.Ceram. Soc., 22, 2427–2434.

55. Gutierrez, C.A., Moreno, R. (2000), ‘Tape casting of non-aqueous silicon nitrideslips’, J. Eur. Ceram. Soc., 20, 1527–1537.

56. Xu, X., Mei, S., Ferreira, J.M.F. (2003), ‘Fabrication of a-SiAlON sheets by tapecasting and pressureless sintering’, Department of Ceramics and Glass Engineering,CICECO, University of Aveiro, Portugal, www.mrs.org/publications/jmr/jmra/2003/jun/012.html.

57. Jung, Y.G., Ha, C.G., Shin, J.H., Hur, S.K., Paik U. (2002), ‘Fabrication of functionallygraded ZrO2/NiCrAlY composites by plasma activated sintering using tape castingand its thermal barrier property’, Mat. Sci. Eng., A323, 110–118.

58. Yeo, J.G., Jung, Y.G., Choi, S.C. (1998), ‘Design and microstructure of ZrO2/SUS316functionally graded materials by tape casting’, Materials Letters, 37, 304–311.

59. Dumont, A.L., Bonnet, J.P., Chartier, T., Ferreira, J.M.F. (2001), ‘MoSi2/Al2O3 FGM:elaboration by tape casting and SHS’, J. Eur. Ceram. Soc., 21, 2353–2360.

60. Chartier, T., Merle, D., Besson, J.L. (1995), ‘Laminar ceramic composites’, J. Eur.Ceram. Soc., 15, 101–107.

61. Carlström, E., Kristoffersson, A. (2002), ‘Water based tape casting and manufacturingof laminated structures’, Key Eng. Mat., 206–213, 205–210.

62. Çalıfl, N. (2004), ‘Functionally graded SiAlON ceramics’, MSc Thesis, AnadoluUniversity, Turkey.

63. Chen, L., Kny, E., Groboth, G. (1998), ‘SiAlON ceramics with gradientmicrostructures’, Surface and Coating Technology, 100–101, 320–323.

64. Jiang, X., Kang, L. (1998), ‘Formation of a-sialon layer on b-sialon and its effect onmechanical properties’, J. Am. Ceram. Soc., 81, 1907–1919.

178

7.1 Introduction

Ceramics have found their use in numerous crosscutting industrial applicationsbecause of excellent hardness, wear, corrosion resistance, and ability towithstand high temperatures. However, ceramics’ reliability and ductilitycompared to metals are not very high. The best approach to increasing thefracture toughness which enables the structural application of ceramics isthrough the development of ceramic composites. Fiber-reinforced compositesdemonstrate the highest fracture toughness and damage tolerance. However,since these materials have a very high density of weak interfaces, they arenot very strong. In addition, their high cost limits their commercial applications.Particulate composites are less expensive to manufacture, but compared tomonolithic ceramics, their fracture toughness increases are insignificant.Several publications on ceramics show that the use of layered materials isthe most promising method for controlling cracks by deflection, bifurcation,microcracking, or internal stresses [1–4]. Layered structures clearly offer akey to greater reliability at a moderate cost and new applications may resultas more complex structures are tailored to specific applications [5].

The way to achieve the highest possible mechanical properties is to controlthe level of residual stresses in individual layers. One can increase the strengthand apparent fracture toughness of ceramics by creating a layer withcompressive stresses on the surface. In such a way, surface cracks will bearrested and, therefore, higher failure stresses are achieved [6]. The variablelayer composition, as well as the system’s geometry, allows the designer tocontrol the magnitude of the residual stresses in such a way that compressivestresses in the outer layers near the surface increase strength, flaw tolerance,fatigue strength, resistance to oxidation, and stress corrosion cracking. In thecase of symmetrical laminates, this can be done by choosing layer compositionssuch that the coefficient of thermal expansion (CTE) of the odd layers issmaller than the CTE of the even ones. The changes in compressive andtensile stresses depend on the mismatch of CTEs, Young’s moduli, as well as

7Design of tough ceramic laminates by

residual stresses control

N O R L O V S K A Y A, Drexel University, USA,M L U G O V Y, Institute for Problems of Materials Science,

Ukraine, J K U E B L E R, EMPA, Lab for High PerformanceCeramics, Switzerland, S Y A R M O L E N K O and

J S A N K A R, North Carolina A&T State University, USA

Design of tough ceramic laminates by residual stresses control 179

on the thickness ratio of layers (even/odd) [7, 8]. The sign and value ofresidual stresses can be established by theoretical prediction [9–12]. Therehave also been a number of experimental studies of laminated ceramics thatwere conducted using these models, attempting to maximize the mechanicalproperties [13, 14].

This chapter consists of the following sections: 7.1 Introduction, 7.2Laminate design for enhanced fracture toughness, 7.3 Processing of Si3N4–TiN and B4C–SiC ceramic laminates, 7.4 Si3N4 based laminates, 7.5 B4Cbased laminates, and 7.6 Future trends. Section 7.2 outlines the main approachto optimizing layered structures leading to an increase of mechanical properties.The technique describes the design of a layered composite with enhancedfracture toughness. The mechanisms responsible for the increase in mechanicalperformance and reliability, acting in the layered ceramics, are consideredhere, too. The main technological steps in manufacturing of ceramic laminatesare considered in Section 7.3. Section 7.4 describes the mechanical performanceof Si3N4 based laminates, and Section 7.5 presents new results on themechanical behavior of B4C based armor laminate materials. A new designwith incremental increase of the apparent fracture toughness of laminates ismentioned in Section 7.6, and a list of references is provided for furtherinformation and reading.

7.2 Laminate design for enhanced fracture

toughness

A number of articles on the design of ceramic laminates leading to a significantincrease of their mechanical properties were published in the past [15–19].Our work is based on the control of thermal residual stresses by optimizationof the layered structure [20, 21]. The proposed approach targets the fracturetoughness increase of laminate ceramic composites and is based on thepreliminary results both from our work [22, 23] and from the work of others[24–26].

7.2.1 Calculation of the residual stresses

A schematic presentation of a general case of two-component multilayeredsample is shown in Fig. 7.1, where ti is the thickness of the ith layer, w is thetotal thickness of the specimen, b is the width, and N is the total number oflayers. Often the two-component brittle layered composites with symmetricmacrostructure are only considered. In this case the layers consisting ofdifferent components alternate one after another and the external layers consistof the same component. The total number of layers N in such a compositesample is odd. Often the layer of each component has some constant thickness,and the layers of same component have identical thickness. In this case all


layers of the first component including two external (top) layers can bedesignated by index 1 (j = 1), and all layers of the second component (internal)can be designated by index 2 (j = 2). The number of layers designated byindex 1 is (N + 1)/2, and the number of layers designated by index 2 is(N – 1)/2.

There are effective residual stresses in the layers of each component in thelayered ceramic composite. During cooling, the difference in deformation,due to the different thermal expansion coefficients of the components, isaccommodated by creep as long as the temperature is high enough. Below acertain temperature, which is called the ‘joining’ temperature, the differentcomponents become bonded together and internal stresses appear. The ‘joining’temperature is difficult to measure experimentally, and in general, it is adoptedto be somewhere below the sintering temperature. In each layer, the totalstrain after sintering is the sum of an elastic component and a thermal component[27, 28]. The residual stresses in the case of a perfectly rigid bonding betweenthe layers of a two-component material are [28]:

s a ar

T TE E f TE f E f1

1 2 2 2 1

1 1 2 2 =

( – ) +

¢ ¢¢ ¢

D(7.1)

and

s a ar

T TE E f TE f E f2

2 1 1 1 2

1 1 2 2 =

( – ) +

¢ ¢¢ ¢

D(7.2)

where ¢E j = Ej/(1 – nj), fN l

w11 =

( + 1)2

, fN l

w22 =

( – 1)2

, Ej and nj are the

x

N

i

2

1

w

xi

ti

y

Surface under tension

7.1 Scheme of a two-component multilayer specimen: numbers oflayers and layer boundary coordinates.


elastic modulus and Poisson’s ratio of the jth component respectively, l1 andl2 are the thickness of layers of the first and second components, aT1 and aT2

are the thermal expansion coefficients (CTE) of the first and second componentsrespectively, DT is the difference in temperature of the ‘joining’ temperatureand current temperature, and w is the total thickness of the specimen. Themismatch of thermal expansion coefficients between different layers inevitablygenerates thermal residual stresses during subsequent cooling of layeredceramics with strong interfaces [28]. The relative thickness of different layersdetermines the relative magnitudes of compressive and tensile stresses, whilethe strain mismatch between the layers dictates the absolute values of theresidual stresses. The important trend is a decreasing of tensile residualstress and an increasing of compressive residual stress with an increase ofthe thickness of layers under tension and a decrease of the thickness of layersunder compression. In this way a change of a layer thickness ratio allowscontrol of the residual stress level in laminates.

7.2.2 Design principles and algorithm

While a number of mechanical properties, such as strength and reliability,can be increased by design of laminates, we specifically target an increase ofthe fracture toughness in our study. In the case of non-homogeneous(particularly layered) materials, the so-called apparent fracture toughnessshould be considered. The apparent fracture toughness is the fracture toughnesscalculated from test data of the layered sample considering this specimen as‘homogeneous’. Such an approach does not meet the fracture mechanicsrequirement of taking into account all features of stress distribution near thecrack tip in layered media, but it is still a useful characteristic allowing aneffective contribution of such factors as residual stresses and materialinhomogeneity to be accounted for. In fracture mechanics, both residual andapplied stresses are usually included in the crack driving force. It can beuseful to consider residual stresses as a part of the crack resistance, andtherefore in laminates with residual compressive stress, the higher resistanceto failure results from a reduction of crack driving force rather than from anincrease in intrinsic material resistance to crack extension [29]. An apparentfracture toughness of laminate can be successfully modeled taking into accounta layered structure. In this case the modeling is a power tool of laminatedesign because it allows prediction of the mechanical behavior of layeredmaterial with a crack.

The compressive residual stress sr1 in the outside layers of a laminateshields natural and artificial cracks in the layer. Therefore, the effective(apparent) fracture toughness of such a structure increases. The morecompressive residual stress induced, the more shielding occurs. Anotherimportant factor that contributes to the apparent fracture toughness increase


is the crack length a. A longer crack promotes more shielding. The maximumlength of a transverse crack in an outside compressive layer is limited by thelayer thickness l1. These two factors determine the apparent fracture toughnessof the material.

In general, the condition for crack growth onset is Ka + Kr = Kc, (Fig. 7.2),where Ka = Ka(sa, a) is the applied stress intensity factor that can be measured,sa is the distribution of applied stress resulted from bending, Kr = Kr(sr, a)is the stress intensity factor due to a residual stress, and Kc is the intrinsicfracture toughness of a material in the layer. If a condition of a crack growthonset is fulfilled then Ka = Kc – Kr is the apparent fracture toughness. If sr1

is compressive, then Kr < 0 and Ka increases. The more | sr1 |, the greater Ka.Similarly, the greater a, the higher Ka. The largest value of a crack length incompressed layer is l1. The maximum apparent fracture toughness can beobtained for such a crack. Unfortunately, small cracks have Ka close to Kc.

A schematic presentation of factors that affect apparent fracture toughnessis shown in Fig. 7.3. The contribution of residual stress to the maximumapparent fracture toughness is Kr = Y(l1/w)sr1 l1

1/2 , where Y(l1/w) is thegeometrical factor [13, 14, 30], and w is the total thickness of a layeredspecimen. The factor Y l w l( / )1 1

1/2 increases as l1 increases (Fig. 7.3(a)). Thecompressive residual stress decreases as l1 increases. It can be calculatedusing equation (7.1). In addition, the residual stress depends on the numberof layers in the sample (Fig. 7.3(b)). The final dependences of Ka on l1 forvarious numbers of layers are shown in Fig. 7.3(c). These dependences arenon-monotonic curves with the maximum that depends on a number oflayers in the laminate. The labels w/5, w/4, w/3 and w/2 designate the maximumthickness of the top layer for symmetrical layered structures with nine, seven,five and three layers, respectively. It can be seen that the highest apparentfracture toughness can be obtained for the three-layer specimen. Thus, thestudy of the layers’ relative thickness and their numbers reveals that themaximum crack shielding will be achieved for three-layer composites withan edge crack extending to the first interface (l1 = w/4) [13, 14]. However,

sa sr sr

saa

7.2 Factors affecting the apparent fracture toughness in a laminate:sa is the distribution of applied stress resulted from bending; sr isthe residual compressive stress; a is the crack length.


the multilayered design is also very important to meet specific requirements,such as ballistic impact, for example. It is essential during impact loading tohave more barriers to arrest cracks. In our case it is the number of compressivelayers. For a three-layer design there is only one such barrier that is a topcompressive layer. The top layer plays a key role, however, multilayereddesign is of further importance to stop cracks more effectively [3, 4].

The design technique to obtain the enhanced fracture toughness of a layeredcomposite is as follows. First, the compositions of the layers are selecteddepending on the intended application of the composite. Then, the relevantmaterial constants entering the design are determined. The constants fordesign are the coefficient of thermal expansion, Young’s modulus, Poisson’sratio, and the density of the corresponding constituents. A very important butat this point of design experimentally unknown parameter is the ‘joining’temperature. Further, effective coefficients of thermal expansion, effectiveYoung’s modulus, average density and the thickness ratio of layers aredetermined using the rule of mixtures.

The next step in the design is the selection of the number of layers. It canbe any appropriate number depending on the required total thickness of thetile. To obtain the enhanced fracture resistance of the layered composite, thefactors affecting the apparent fracture toughness should be taken into account.Usually, the thickness of the thinnest possible layer is limited by the

(a)

(b)

(c)

w/5 w/4 w/3 w/2 l1

sr

Ka

Kc

Y(l1/w) l11/2

9 layers

7 layers

5 layers

3 layers

3 layers

7.3 Factors affecting laminate design for maximum apparent fracturetoughness.


manufacturing technology. Note that a compressive layer should be thinenough to reach a high level of residual stress. Another important requirementis the thickness ratio of layers with high CTE (tensile stress) and layers withlower CTE (compressive stress). Any appropriate thickness ratio can be usedas a first approximation. Then the tensile layer thickness is found. After this,the residual stresses are calculated using equation (7.1) and (7.2). The totalthickness of the sample is also determined at this step for a given layer’sthickness ratio taking into account the selected number of layers. The thicknessratio is changed after the analysis of residual stress and the total thickness ofthe specimen. Note that increasing the ratio of tensile layer thickness tocompressive layer thickness decreases tensile residual stress. However, itcan result in increasing the total thickness of the sample. After changing thethickness ratio, the calculation is repeated. Such iterations are continued tofind a unique optimal layer thickness ratio that produces the maximum possiblecompressive residual stress, low tensile residual stress, and required totalthickness of the sample.

A design algorithm is presented in Fig. 7.4. The maximum possible apparentfracture toughness of the corresponding layered structure is also determinedin all iterations as an indicative parameter of the design. The determinationof the apparent fracture toughness uses the compressive residual stress andthe thickness of an outside layer as a crack length at any given iteration.These two parameters (the compressive residual stress and the thickness of

Selection of compositionInput parameters:moduli, CTEs, etc.

Layer effective properties calculation

Selection of layer number andthickness of layer under compression

Thickness ratio and tensile layerthickness determination

Residual stress calculation

Analysis of the residual stress andthickness ratio changing

Output parameters:layer thicknesses

7.4 An algorithm of laminate design.


the top layer) have trends acting in opposite directions. A decrease in the toplayer thickness can increase the residual stress in the layer, but it decreasesthe length of the maximum crack. Therefore, the maximum apparent fracturetoughness was always used to analyze the optimal thickness ratio. It wasshown also that in order to obtain the higher resistance to failure, the tensilelayer should be made as stiff as possible (i.e., high elastic modulus), whereasthe compressive layers should be as compliant as feasible (i.e., low elasticmodulus) [11].

7.2.3 Calculations of the apparent fracture toughness

A weight function analysis has been used to estimate the apparent fracturetoughness in laminates with residual stresses [10, 21, 31, 32]. A schematicpresentation of the analyzed crack location in the layered specimen is presentedin Fig. 7.5, where a is the crack length and n is the number of layers crossedby the crack. The choice of coordinate system is of great importance to theapparent fracture toughness calculations because of a significant simplificationof the procedure. The most appropriate coordinate origin is on the tensilesurface of the sample under bending. The geometry of the multilayeredmaterial analyzed here is such that the problem can be reduced to one dimensionand that analytically tractable solutions can be used [21].

An experimental value of the apparent fracture toughness can be foundusing the expression [10, 21]:

Ka = Y(a)sma1/2 (7.3)

where Y ( ) = 1.99 – (1 – ) (2.15 – 3.93 + 2.7 )

(1 + 2 )(1 – ),

2

3/2a a a a aa a

s mP s s

bw =

1.5 ( – )1 22

and

y

Surface under tension

a

w

x

N

n + 1

n

2

1

7.5 An analyzed crack location in a layered sample.


a = a/w, P is the critical load (the applied bending load corresponding tospecimen failure) and s1 and s2 are outer and inner support spans of the four-point bending fixture.

The apparent fracture toughness of a layered composite can be calculatedanalytically by [10, 21]:

Ka

= 6 ( ) ( – )( – )

, [ – ] + , [ – ]

1/21

20 2 1

( )

2+1 0 1 =1 0 1

–1

Y a I I I K K

w E h xa

I x I dx E h xa

I x I dx

L L L ci

r

nx

a

L L i

n

ix

x

L L

n i

i

a

a a¢ ÊË

ˆ¯ ¢ Ê

Ëˆ¯

ÏÌÔ Ú ÚSÓÓÔ

¸˝ÔÔ

(7.4)

where K ci

1( ) is the intrinsic fracture toughness of the ith layer material, Kr is

the stress intensity due to the residual stresses, h(x/a, a) is the weight functionfor an edge-cracked sample [9, 19, 30], xi is the coordinate of an upperboundary of the ith layer (Fig. 7.1), ¢Ei = Ei/(1 – ni), and Ei and ni are theelastic modulus and Poisson’s ratio of the ith layer, respectively. The expressionsfor ILj (j = 0, 1, 2) and JLj (j = 0, 1) were obtained from Ref. [21] as follows:

Ij

E x xLj i

N

i ij

ij = 1

+ 1 ( – )

=1

+1–1+1S ¢ (7.5)

Jj

E x xL j i

N

i i ij

ij = 1

+ 1 ( – )

=1

+1–1+1S e ¢ (7.6)

where e i is the strain in the ith layer, which is not associated with any stress.The thermal expansion and/or volume change due to a crystallographic phasetransformation might be the source of this strain. However, the case of aphase transformation is beyond the scope of this chapter. In the case ofthermal expansion:

e biT

T

i

j

T dT = ( )0Ú

where bi(T) is the thermal expansion coefficient of the ith layer at thetemperature T. T0 and Tj are the actual and ‘joining’ temperatures, respectively.If bi(T) is a linear function, e bi i T > ,£ D where DT = Tj – T0,

< ( ) + ( )

20b

b bi

i i jT T≥ is the average value of the thermal expansion

coefficient in the temperature range from T0 to Tj.The stress intensity due to residual stresses is given by equation (7.7) [10,

21]:


KI I Ir

L L L

= 1 – 1

20 2

¥ ¢ÏÌÓ

ÊË

ˆ¯Ú , [ – + ( – ) ] ++1 1 1 2 0 1 0 0 1E h x

aI J I J I J I J x dxn

x

a

L L L L L L L Ln

a

+ , [ – + ( – ) ]=1 1 1 2 0 1 0 0 1

–1

Si

n

ix

x

L L L L L L L LE h xa

I J I J I J I J x dxi

i

¢ ÊË

ˆ¯

¸˝˛Ú a

(7.7)

The apparent fracture toughness Ka in layered specimens can be analyzedas a function of the crack length parameter a, where a Y a = ( ) .1/2a Thecrack length parameter a is the most appropriate to demonstrate criticalconditions of crack growth. One of the advantages of this parameter is thatthe stress intensity factor of an edge crack for a fixed value of the appliedstress sm is a straight line from the coordinate origin in the coordinatesystem Ka – a . Since K am1 = ,s ˜ the slope of the straight line is the appliedstress sm. The conditions for unstable crack growth in the internal stressfield are as follows [10, 21]: K1(sm, a) = Ka(a); dK1(sm, a)/da ≥ dKa(a)/da.Using parameter a , these conditions become s m aa K a˜ ˜ = ( ) ands m adK a da ( )/ ,≥ ˜ ˜ which can be reduced to:

K a a dK a daa a( )/ ( )/˜ ˜ ˜ ˜≥ (7.8)

It follows from equation (7.8) that unstable crack growth occurs if theslope of the straight line corresponding to the stress intensity factor at constantapplied stress is greater than or equal to the slope of the tangent line to thefracture resistance curve at the same point (Fig. 7.6). Also the applied stressintensity factor becomes higher than the fracture resistance of the material.

7.2.4 Other toughening mechanisms in laminates

In addition to a crack shielding phenomenon that exists due to residual stressthere are two other crack deflection mechanisms leading to a laminatetoughening. Cracks that form in one layer can be deflected either along weakinterfaces with adjacent layers [33, 34] or into layers with compressiveresidual stresses [35, 36]. Since we investigated laminates with strong interfaces,in our design the cracks were not deflected along interfaces; however, crackkinking and bifurcation in layers under compression has often been observedduring mechanical testing (Fig. 7.7). It was shown that the crack tends todeviate as it approaches the centerline of layers with compressive stresses[7]. When the crack enters the layer with compressive stresses, tensile stressesappear near the edges of the crack. These tensile stresses are parallel to the


free surfaces (crack moving direction). They appear because compressivestresses, which are perpendicular to the free surfaces, have to become zeroon the free surface. These tensile stresses are maximal at the centerline oflayers. The condition of kinking (and bifurcation) of a crack along the centerlineis

s r cl AK12

12 >

where A is a constant and Kc is the fracture toughness of layers with compressivestresses [7]. The constant A takes different values for kinking and bifurcation.When compressive stresses are very high or the fracture toughness of layersis sufficiently low, the crack not only deviates at the centerline, but alsobifurcates, preventing the catastrophic failure of the sample during the bendingtest. Compressive stresses, as well as thickness of layers, influence bothcrack deviation and bifurcation behavior [8, 37–39]. It was shown that arange of layer thicknesses exists where crack bifurcation can occur [40].

Ka /ã = sm

dKa/dã

ã0

Kapp Tangentline

Fractureresistance

Stress intensity factorat constant applied

stress

7.6 General criterion of stable/unstable crack growth in a brittlematerial.

9 mm

7.7 Crack bifurcation in the B4C–SiC woven fabric based laminate.The B4C and SiC woven fabric layers cannot be distinguished underthe optical microscope.


This range depends on the elastic constants of the layers, the layer number,the mismatch in CTEs, and DT.

The important case of specimens with a fixed total thickness was consideredin Ref. [40]. There are certain features of crack bifurcation under theseconditions, such as that if the sample with a fixed total thickness has toolarge a number of layers there will be no bifurcation. Layer thickness andcomposition are important and efficient parameters to control the bifurcationin laminates. The effect is comparable with a crack bridging phenomenon[21]. The bifurcation mechanism increases the laminate fracture toughnessby approximately 1.5–2 times.

7.3 Processing of Si3N4–TiN and B4C–SiC ceramic

laminates

Two systems of ceramic laminates have been chosen for manufacturing andmechanical testing of laminates – Si3N4 and B4C based ceramics. They arewell-known materials for crosscutting industrial applications that can beused as cutting tools, igniters, wear parts, armor, etc.

Silicon nitride is one of the most promising ceramics for structuralapplication, with good corrosion resistance and outstanding mechanicalproperties. Residual stresses can be created in dense Si3N4 material byincorporating a dispersion of metal-like refractory compounds, such as nitridesor carbides. Such TiN particles create compressive tangential stresses in thesurrounding matrix while remaining under tensile stress. From the point ofview of chemical compatibility TiN is the most promising dispersion becauseof its stability in contact with Si3N4 under sintering conditions [41]. It hasbeen demonstrated that the fracture toughness and strength of silicon nitridecan be increased in composites [42]. The addition of TiN leads to increasesin Young’s modulus, CTE, electrical conductivity, etc. [43]. At the sametime, some properties of silicon nitride are reduced by these inclusions.Usually the metal-like refractory compounds have a much lower oxidationresistance than that of silicon nitride. Their additions result in a decrease inthe oxidation resistance of the initial matrix material [44]. Also, these additionsaccelerate the high-temperature creep rate, which increases drastically whenthe TiN content is higher than 30 wt%.

Boron carbide is another important ceramic material with many usefulphysical and chemical properties. After cubic boron nitride, it is the hardestboron-containing compound [45]. Its high melting point, high elastic modulus,large neutron capture section, low density, and chemical inertness makeboron carbide a strong candidate for several high-technology applications.Due to its low density and superior hardness, boron carbide is a very promisingmaterial for lightweight ballistic protection. Boron carbide exists as a stablesingle phase in a large homogeneity range from B4C to B10.4C [46]. The most


stable boron carbide structure is rhombohedral with a stoichiometry of B13C2,B12C3, and some other phases close to B12C3 [47, 48]. The Vickers hardnessof B4C is in the range of 32–35 GPa [49]. There is an indication that thehardness of stoichiometric B4C is highest in comparison with that of boron-rich or carbon-rich boron carbide compositions [50–52]. However, B4C-based composites have a relatively low fracture toughness of 2.8–3.3 MPam1/2 [53]. While high hardness is one of the very important requisite indicatorsfor a material’s ballistic potential, toughness might play an equally importantrole in realizing that potential. Thus, materials with both high hardness andhigh fracture toughness are expected to yield the best ballistic performance[54]. Therefore, a significant increase in fracture toughness of boron carbide-based laminates has the potential for realization of improved armor materialsystems.

Typically the manufacturing steps of ceramic laminates include (a) ballmilling of powders in certain proportions; (b) rolling of thin tapes; (c) stackingof rolled tapes; and (d) hot pressing of the stack. A schematic presentation ofthe manufacturing steps is shown in Fig. 7.8. This process was used tofabricate layered specimens of the Si3N4–TiN and the B4C–SiC systems. ForSi3N4 based ceramic laminates, powder mixtures of the following compositionswere used:

∑ Silicon nitride (95% of a-modification) with 2 wt% of alumina and 5 wt%of yttria as sintering aids (layers with index 1, bulk compressive stress)

∑ The above mixture with the addition of 20–50 wt% TiN or pure TiN(layers with index 2, bulk tensile stress).

a-SiC and TiN powders with grain size 1 mm and 3 mm respectively wereused for layered sample fabrication. For B4C based ceramic laminates, powdermixtures of the following compositions were used:

Grinding Sieving Plasticization

Rolling Stacking Hot pressing

Crude rubber + petrol

7.8 Schematic presentation of manufacturing steps of ceramiclaminates.


∑ Mixture of boron carbide with 30 wt% SiC (layers with index 1, bulkcompressive stress)

∑ Pure boron carbide (layers with index 2, bulk tensile stress).B4C and a-SiC powders with a grain size of 2–5 mm were used forlaminate manufacturing.

A detailed description of the manufacturing process is provided below.The mixtures of various compositions were milled in the ball mill for 48 h.The average grain size of the milled powders was about 1 mm. Crude rubber(4 wt%) was added to the mixture of powders as a plasticizer through a 3%solution in petrol. The powders were then dried up to 2 wt% residualamount of petrol in the mixture. After sieving powders with a 500 mm sieve,granulated powders were dried up to 0.5 wt% residual petrol. A roll mill with40 mm rolls was used for rolling. The velocity of rolling was 1.5 m/min. Theworking pressure was about 10 MPa to obtain a relative tape density of64%. The thickness of green tapes was 0.4–0.5 mm and the width was up to100 mm.

The formation of a thin ceramic layer is of specific importance, as thesizes of residual stress zones (tensile and compressive) are directly connectedto the thickness of layers. The advantage of rolling, as a method of greenlayer production, is that it allows easy thickness control, achieves high greendensity of the tapes, and requires a rather low amount of solvent and organicadditives compared to other methods such as tape casting [55]. Additionalpowder refinement, giving a higher sintering reactivity, might occur due tothe large forces applied in the pressing zone during rolling. The modeling ofrolling, as recently performed, potentially allows optimizing the process ofroll compaction [56]. There is a challenging problem to produce thin tapeswith a small amount of plasticizer and sufficient strength and elasticity tohandle green layers after rolling.

A schematic presentation of rolling is shown in Fig. 7.9. Powders arecontinuously supplied in the bunker and further into the deformation zone inbetween rolls. Powders are supplied to the deformation zone due to both thegravitational force and friction between rolls and powders. The relative densityof the tape (rr) can be calculated from

rrl

ar

p

s

Rh

= 1 + 2Ê

ËÁˆ¯

(7.9)

where rp is the relative powder density, l is a drawing coefficient, a is theintake angle, and R is the roll diameter.

Green tapes were stacked together to form the desired layered structuresand ceramic samples were prepared by hot pressing the stacked tapes. Thehot pressing was performed at 1820oC and 30 MPa for 45 min for Si3N4

based laminates [22], and at 2150–2200oC and 30 MPa for 50–60 minutes


for B4C based laminates [57]. Graphite dies were used for the hot pressingof laminates with graphite surfaces coated by a BN layer in order to preventdirect contact between graphite and ceramic material. During hot pressing oflaminates, shrinkage of the individual layers occurred, and their thicknessreduced to 0.15 mm after hot pressing. The interfaces between individuallayers of the same composition completely disappeared and only the interfacebetween layers of different compositions could be distinguished. Dense (95–100% density) laminate samples were obtained. The specimens for mechanicaltests were prepared by machining them from the hot pressed tiles. StandardMOR bars of dimension 50 mm ¥ 4 mm ¥ 3 mm were surface ground to thespecification stated in EN843-1. The bars were also chamfered along thelong edges with a chamfer angle at 45o to a dimension of 0.12 ± 0.03 mm.The fracture toughness was measured by the SEVNB technique [58, 59]using equation (7.3). V-notches with tip radii of the order of 10–15 mm weremade in the specimens by a diamond saw followed by notching with a razor

100 mm

(b)

B4C tapesB4C–30 wt% SiC tapes

1 Bunker, 2 Powders, 3 Rolls, 4 Transmission, 5 Motor,6 Bottom support, 7 Tape

(a)

1 2 3 4 5

67

w1w2w2

V2

V1

7.9 (a) Schematic presentation of rolling; (b) a photograph of B4C andB4C–30wt%SiC rolled tapes. The thickness of an individual tape afterrolling is between 0.4 and 0.5 mm.


blade with diamond abrasive to obtain a sharp notch tip. The elastic moduluswas measured by a standard four-point bending technique.

7.4 Si3N4 based laminates

7.4.1 Mechanical properties

Four different laminate composites, Si3N4/Si3N4–20wt%TiN, Si3N4/2(Si3N4–20wt%TiN), Si3N4/Si3N4–50wt%TiN and Si3N4/TiN, were chosen to studytheir mechanical performance. The parameters of their components, such asCTE and Young’s modulus (compiled from literature data), composition andlayer thickness, are given in Tables 7.1 and 7.2. The expression 2(Si3N4–20wt%TiN) means that the (Si3N4–20wt%TiN) layer is twice as thick as theSi3N4 layer (see Table 7.2). Besides these four designs, one more design ofSi3N4/Si3N4 laminate was used as a base for comparison. The laminates ofthis design were prepared in the same way as the others, though all layerswere of the same composition. Therefore, no residual stresses can appearduring cooling. It is worth noting that both the Young’s modulus and fracturetoughness of these Si3N4/Si3N4 laminates were measured to be on the samelevel as those of standard Si3N4 ceramics prepared by the standard powderroute, which includes no rolling. The strength of the Si3N4/Si3N4 laminatewas less than that of the standard Si3N4 ceramics with values of 508 ± 3.2and 750 ± 20.7 MPa, respectively. Mechanical properties such as the strength,Young’s modulus, and fracture toughness of the laminates are presentedin Table 7.3. As one can see from Table 7.3, while the strength of

Table 7.1 Young’s moduli and CTE of the components

Composition E (GPa) CTE (10–6 K–1)

Si3N4–5wt%Y2O3–2wt%Al2O3 320 3Si3N4(5wt%Y2O3–2wt%Al2O3)–20wt%TiN 335.62 3.826Si3N4(5wt%Y2O3–2wt%Al2O3)–50wt%TiN 364.93 5.378TiN 440 9.35

Table 7.2 Calculated residual stresses in Si3N4 based laminates

Composition Thickness of layers scomp stens

(mm) (MPa) (MPa)

Si3N4 Si3N4 withTiN

Si3N4/Si3N4–20wt%TiN 250 210 188 247Si3N4/2(Si3N4–20wt%TiN) 245 530 280 151Si3N4/Si3N4–50wt%TiN 200 330 765 516Si3N4/TiN 200 400 2467 1078


Si3N4/Si3N4–20wt%TiN laminates is approximately on the same level asthat of the Si3N4/Si3N4 laminates, further increases of the TiN content to50% and 100% resulted in a significant decrease of both strength and Young’smodulus. The measured fracture toughness of the Si3N4/TiN laminates alsoshowed a decrease similar to strength and Young’s modulus values. Anexplanation is sought for this reduction in mechanical properties.

The Si3N4/Si3N4–20wt%TiN laminates showed an increase in apparentfracture toughness. This increase can be explained by the introduction of theresidual bulk compressive stresses in the Si3N4 layers. In the case where thethicknesses of the Si3N4 and the Si3N4–20wt%TiN layers were similar, thecalculated residual compressive stress was about 188 MPa and the residualtensile stress about 246.5 MPa. The measured value of the apparent fracturetoughness was 7.41 ± 1.79 MPa m1/2. There was a further increase in KIC (8.5± 0.01 MPa m1/2) for the laminates with 20 wt% TiN when the relativethickness of the Si3N4–20wt%TiN layers was increased compared to thethickness of the pure Si3N4 layers. The reason for this is a significant increaseof the residual compressive stress, and at the same time a decrease of theresidual stress in the Si3N4–20wt%TiN layers (Table 7.2). However, an increaseof TiN content to 50 wt% resulted in a significant increase of the residualtensile stress in the laminates. The calculated tensile stress values are higherthan the tensile strength of the material, and there is therefore much crackingand a decrease in all mechanical properties (Table 7.3).

7.4.2 Apparent fracture toughness of the layeredcomposite with residual compressive or tensilestresses in the top layer

A detailed study of the effect of the residual compressive or tensile stressesin the top laminate layers on the apparent fracture toughness values has beendone for Si3N4/Si3N4–30wt%TiN. Young’s moduli of the Si3N4 and the Si3N4–30wt%TiN monolithic samples, used as reference materials, were measuredto be 308 GPa and 323 GPa, respectively. Mean values of intrinsic fracturetoughness of monolith materials, measured by SEVNB, are approximately

Table 7.3 Mechanical properties of Si3N4 based laminates with differentlayer compositions

Composition sf (MPa) E (GPa) KIC (MPa.m1/2)

Si3N4/Si3N4 508 ± 3 307 5.54 ± 0.01Si3N4/Si3N4–20wt%TiN 356 ± 76 313 7.41 ± 1.79Si3N4/2(Si3N4–20wt%TiN) 450 ± 83 — 8.50 ± 0.01Si3N4/Si3N4–50wt%TiN 158 ± 15 298 —Si3N4/TiN 141 ± 11 157 3.97 ± 0.52


the same both for the Si3N4 and for the Si3N4–30wt%TiN compositions andare equal to 4 ± 1 MPa m1/2. These measured values of Ei and K IC

i( ) were usedin all following calculations. The calculated values of the apparent fracturetoughness as a function of the crack length parameter a in the Si3N4/Si3N4–30wt%TiN laminate with compressive outer layers are shown in Fig. 7.10(a).The toughness increases in the layers with compressive stress with increasingcrack length, and it decreases in the layers with tensile stress as the crackcontinues to grow. The layers with compressive and tensile stresses areshown in Fig. 7.10 in white and gray colors, respectively. As one can see, Ka

reaches its maximum or minimum values as the crack approaches the interfacewith a new layer of an opposite stress sign. For the first Si3N4 top layer withcompressive stress, the calculated apparent fracture toughness increases from3.9 to 17 MPa m1/2 as a function of the crack length parameter. Theexperimentally measured Ka values, presented as solid circles in Fig. 7.10(a),show an excellent fit with the calculated values. The crack length parametersfor the experimentally measured Ka were calculated from the initial notchlengths. All experimentally measured points are located on close to a straightline between the coordinate origin and the maximum Ka point at the interfacebetween the first and second layers. Failure of all samples occurred at 351 ±13 MPa. The calculated Ka decreases in the second Si3N4–30wt%TiN layerwith a residual tensile stress from 17 to 5 MPa m1/2, followed by the nextincrease from 5 to 14 MPa m1/2 in the third Si3N4 layer with a residualcompressive stress. The insert in Fig. 7.10(a) shows an optical micrograph of

0 0.02 0.04 0.06 0.08 0.10

(a)ã (m1/2)

Si3N4 + 30% TiNSi3N4

16

12

8

4

0

Kapp (MPa.m1/2)Si3N4

Si3N4-30% TiN Si3N4 Si3N4-30% TiN

(b)

0 0.02 0.04 0.06 0.08 0.10ã (m1/2)

Kapp (MPa·m1/2)

Si3N4Si3N4 + 30% TiN

8

6

4

2

0Si3N4-30% TiN Si3N4 Si3N4-30% TiN Si3N4

7.10 Apparent fracture toughness as a function of crack lengthparameter a in the laminate with compressive (a) and tensile (b)outer layers. Filled circles correspond to the experimental data.Inserts are optical micrographs of the two parts of Si3N4/Si3N4–30wt%TiN laminate samples with (a) Si3N4 surface layers with aresidual compressive stress and (b) Si3N4–30wt%TiN surface layerswith a residual tensile stress after single-edge V-notch beam test.


two parts of the Si3N4/Si3N4–30wt%TiN laminate sample with a V-notch inthe top layer with residual compressive stress after the SEVNB test. As onecan see, there is a relatively straight crack path with no sharp crack deviation,deflection, or bifurcation during the crack propagation.

Figure 7.10(b) shows the calculated apparent fracture toughness as a functionof the crack length parameter a in the Si3N4/Si3N4–30wt%TiN laminatewith a residual tensile stress in the outer layers. The toughness decreasesfrom 3.9 to 0.8 MPa m1/2 within the first Si3N4–30wt%TiN layer as the crackreaches the first interface. Toughness increases from 0.8 to 6.4 MPa m1/2 inthe second Si3N4 layer with a residual compressive stress, and it decreasesagain from 6.4 to 1 MPa m1/2 within the third Si3N4–30wt%TiN layer witha residual tensile stress. There is no continuous growth of the crack in thiscase. The crack starts to propagate and then becomes arrested; after this itcontinues to grow again. The crack arrest results in a ‘pop-in’ event on theload–displacement diagram (Fig. 7.11). A stress of such ‘pop-in’ event is theonset stress of crack propagation. This stress, as well as an initial notchlength, was used to calculate the measured apparent fracture toughness.Experimentally measured values of Ka fit well with the calculated numbers.The experimental data can be considered in two different sets. The first setincludes Ka measured with notch tips within the first Si3N4–30wt%TiN andthe second Si3N4 layers. Failure of all samples from the first set occurred at116 ± 2 MPa. The second set includes two Ka values measured with notchtips within the third Si3N4–30wt%TiN layer. Failure of these two samplesoccurred at 71 ± 1 MPa. The insert in Fig. 7.10(b) shows an optical micrographof two parts of the Si3N4/Si3N4–30wt%TiN laminate sample with the V-notch placed in the Si3N4–30wt%TiN top layer with a residual tensile stress

Final load

Pop-in-load

0 10 20 30 40 50 60 70 80 90Deflection (mm)

200

180

160

140

120

100

80

60

40

20

0

Load

(N

)

7.11 Load–displacement diagram of single-edge V-notch beamsample with pop-in.


after the SEVNB test. As one can see from the optical image, the crack pathdeviates strongly from a straight line with 90o crack deflection occurring inthe center of each Si3N4 layer with a residual compressive stress. Whiletraveling only a short distance of about one Si3N4–30wt%TiN layer thicknessalong a centerline, the crack kinks out into the Si3N4–30wt%TiN layer witha residual tensile stress.

The calculations indicate an unambiguous trend for the apparent fracturetoughness behavior. The value of Ka increases in the layers with residualcompressive stress and decreases in the layers with residual tensile stress asa function of the crack length (or the crack length parameter). The calculatedincrease of Ka is confirmed by the experimental data in the laminates withthe compressive outer layer (Fig. 7.10(a)). As one can see from Fig. 7.12(a),cracks with crack length parameter from 0 to point A1 will demonstrateunstable crack growth. In this case, once the crack starts to propagate at acertain stress, it cannot be arrested; this results in complete failure of thesample, since the applied stress intensity factor is always higher than thefracture resistance of the laminate. Cracks with crack length parameter betweenA1 and A3 propagate in two stages. For example, a crack with crack lengthparameter at A2 will have an unstable growth from point B2 to point C on theKapp– a plot (Fig. 7.12(a)). Stable growth of this crack will occur from pointC to point D. For all cracks with crack length parameter a from A1 to A4,failure occurs at a stress equal to the slope of the straight line OD, which isa threshold stress. The threshold stress sthr is determined by the maximumvalue of Ka at the interface between the first (compressive) and the second

Stress intensity factor atconstant applied stress

Compressivelayer

Tensilelayer

Compressivelayer

Thresholdstress

ã0

A1A2 A3 A4

Kapp

KmD

C

B1

B2B3

Kc

(a) (b)

Stress intensityfactor at

constant appliedstress

0

Kapp

Km

Kc

A1 A2 A A3 ã

Compressivelayer

Tensilelayer

Tensilelayer

Thresholdstress sthr

D

C

B1 B2

BB3

B¢

B¢3

s0B

s0B1

7.12 Conditions for stable/unstable crack growth in a layeredstructure: (a) a range of crack length parameters for stable crackgrowth in a laminate with a residual compressive stress in a toplayer; (b) stable/unstable crack growth in a laminate with a residualtensile stress in a top layer.


(tensile) layers, and no failure can occur below sthr if the sample contains thesurface cracks located only in the first layer. The curvature of the Ka plot isa function of a value of the residual stress. The higher the residual stress, themore concave the curvature of Ka. At a certain small value of residualcompressive stress, the line OD can have only one intersection point with theKa plot, and therefore no stable crack growth stage can occur.

The conditions for stable/unstable crack growth in the laminate with residualtensile stress in the top layer are shown in Fig. 7.12(b). Cracks with cracklength parameter A1 for such laminates will propagate only unstably at stresslevels above sOB1. Cracks with crack length parameter A grow unstably atstress sOB. This unstable growth occurs between points B and C (Fig. 7.12(b)),because the points belonging to the BC segment lie above the Ka plot. Atpoint C, the condition of equation (7.8) is violated and the crack growthbecomes stable between points C and D, which means that any crackadvancement requires an increase of the applied stress. Point D is a maximumvalue of Ka at the interface between the second (compressive) and the third(tensile) layers. This point determines a stress s0D = sthr. Above s0D, thecrack propagates unstably until complete failure. In such a way all initialcracks in the first (tensile) and the second (compressive) layers with a cracklength parameter greater than A2 (Fig. 7.12(b)) will initiate specimen failureat the same s0D = sthr stress value. The initial cracks with tips in the third andfourth layers will initiate specimen failure at the different stress value that isdetermined by the maximum value of Ka at the interface between the fourthand fifth layers. This stress is sthr for cracks with tips located in the third andfourth layers. It should be noted that points B and B3 in Fig. 7.12(b) correspondto the measured Ka values (using ‘pop-in’ stress), while points B¢ and ¢B3 orbelonging to the 0D straight line are determined by the initial notch lengthand the failure stress of the sample.

As implied by the above analysis, the surface cracks which have sufficientlength to fall into the region of stable crack growth will all cause failure atthe same sthr stress. At the same time, if a residual compressive stress in thetop layer is not high enough, the small cracks can cause catastrophic failureonce they start to grow. Therefore, it might be that different mechanismssuch as crack bridging or transformation toughening can be more effective inpreventing small cracks from growing unstably.

As a result of this part of the work, the apparent fracture toughness as afunction of the crack length parameter a Y a = ( ) 1/2a has been calculated forthe Si3N4/Si3N4–30wt%TiN laminates with residual compressive or tensilestresses in the top layers. The toughness increases in the layers with compressivestress as the crack length increases, and it decreases in the layers with tensilestress as the crack continues to grow. The experimentally measured Ka valuesfor the laminates show an excellent fit with the calculated values. It wasfound that a threshold stress exists for cracks of a certain length. Stable crack


growth occurs for the majority of cracks with a threshold stress as indicatedby the K aa – ˜ graph. Short cracks will propagate unstably because theapplied stress intensity factor is always higher than the fracture resistance ofthe laminate.

If the residual compressive stress is small enough, a situation can occurwhere no stable crack growth exists for cracks with tips located within thefirst compressive layer. Therefore, it is important to introduce high residualcompressive stresses that will provide a steep slope of the apparent fracturetoughness curve to include short cracks in the region of stable crack growth.Obtaining a high residual compressive stress in the first layer is an effectiveway of providing high toughness at small crack lengths, thereby ensuringimproved flaw tolerance and surface damage resistance.

7.4.3 Fracture surfaces after fracture toughness tests

The typical fracture surfaces of the pure Si3N4 layer and the Si3N4–20wt%TiNlayer are shown in Fig. 7.13. The bimodal grain size distribution exists witha number of elongated grains being surrounded by small rounded grains ofSi3N4. The average grain size in the Si3N4 layer was 0.4–0.5 mm. Themicrograph of the Si3N4–20wt%TiN fracture surface revealed that a majorityof the grain sizes were in the range of 1–2 mm, with some grains of size lessthan 1 mm. It was shown that the TiN has a homogeneous distribution in theSi3N4 matrix and no solid solution was detected between Si3N4 and TiNparticles [60].

Fracture surfaces of Si3N4/Si3N4, Si3N4/Si3N4–20wt%TiN, Si3N4/2(Si3N4–20wt%TiN) and Si3N4/TiN laminates are shown in Fig. 7.14. The fracturesurface of the Si3N4/Si3N4 laminate, where no residual stresses were generatedduring cooling, is flat and smooth (Fig. 7.14(a)). As layers of a different

(a) (b)

7.13 Micrographs of fracture surfaces of (a) Si3N4 layer, and (b)Si3N4–20wt%TiN layer, in Si3N4/Si3N4–20wt%TiN laminate.


composition are used, the fracture surface becomes rougher. For the Si3N4/Si3N4–20wt%TiN laminate, there are two zones on the fracture surface. Thefirst zone near the notch tip has a rough surface and corresponds to a slowcrack growth. The second zone has a rather smooth surface with distinctsteps only at the interfaces between layers. This zone corresponds to a fastcrack growth (Fig. 7.14(b)). No crack bifurcation occurred and two equalparts of the sample could be found after failure. The Si3N4/2(Si3N4–20wt%TiN)laminates failed after crack bifurcation. The part of the fracture surface nearthe notch tip was the same as those shown in Fig. 7.13. At the moment whenthe crack bifurcated, an unusually smooth fracture surface was observed(Fig. 7.14(c)). When the value of residual tensile stresses approaches thevalue of the tensile strength of the layer, cracks in the layers are generated,as was the case in the Si3N4/Si3N4–50wt%TiN and Si3N4/TiN laminates. The

(a) (b)

(c) (d)

7.14 Fracture surface of laminate composite: (a) Si3N4/Si3N4laminates; (b) Si3N4/Si3N4–20wt%TiN laminates; (c) Si3N4/2(Si3N4–20wt% TiN) laminates and (d) Si3N4/TiN laminates.


cracks originated during the cooling stage after the hot pressing of the laminatesand appeared due to the large mismatch of CTEs and elastic moduli. Channelcracks were observed in the laminates with a difference in compositionbetween layers, starting with 50 wt% TiN content and higher. Si3N4/TiNlaminates demonstrate channel cracking (Fig. 7.14(d)) similar to the cracksdescribed in Ref. [24]. These cracks are responsible for the dramatic decreasein the mechanical properties of Si3N4 based laminates. To reduce or eliminatecracking, it is necessary to make composites with more similar characteristicsbetween the layers, especially the CTE and elastic moduli. The extent ofchannel cracking was decreased in laminates with Si3N4–50wt%TiN layersin comparison to composites where one of the layers was pure TiN. Channelcracking was fully eliminated for composites with a Si3N4–20wt%TiN layercomposition. An absence of pre-existing cracks resulted in an increase of thestrength and fracture toughness.

The fracture surface of Si3N4/Si3N4–50wt% TiN is shown in Fig. 7.15. Asone can see, there is a high roughness of the surface, and bifurcation of themoving crack occurred when it was inside the Si3N4 layer with residualcompressive stresses. There are fracture steps and channel cracks at theSi3N4–50wt%TiN layers which are perpendicular to the interfaces of thecomposite. The fracture steps appeared only at layers with tensile stresses.Such fracture steps and other defects are responsible for a decrease inmechanical properties. Multiple bifurcations occur for pre-existing cracksinside the layers with residual compressive stresses, and in addition themoving crack bifurcates during sample loading.

7.5 B4C based laminates

7.5.1 Design and mechanical behavior

The material systems selected for the proposed study were B4C and B4C–30wt%SiC because of their promise for ballistic applications [61–63]. Asymmetric three-layered composite was considered for the design andmanufacture of armor tiles as shown in Fig. 7.16 [64]. Table 7.4 shows therelevant material constants entering the design (compiled from the literature),and Table 7.5 shows the corresponding calculated residual stresses in thethree layered B4C/B4C–30wt%SiC laminate. The maximum possible apparentfracture toughness for corresponding layered structures is also presented.The layers under tensile stress have higher CTE, and in this case they areB4C layers. The layers under compressive stress have lower CTE; here theyare B4C–30wt%SiC layers. A temperature T = 2150∞C was used for themajority of the calculations, when residual stresses appeared in the layersupon cooling from the hot pressing temperature. There is no liquid phasepresent during the sintering of B4C/B4C–SiC ceramics [65], therefore the hot


pressing temperature was used as the ‘joining’ temperature DT for calculations.All laminates were designed in such a way that the tensile stresses had beenmaintained at low values. It should be noted that the value of KIC given inTable 7.2 is a theoretical maximal apparent fracture toughness that was usedto estimate the maximum possible toughening of the three-layered laminate.The experimental values measured using the single-edge V–notch beam(SEVNB) method yielded 7.42 ± 0.82 MPa m1/2 [66], which is still a veryhigh value for brittle boron carbide based composites. Thus, the proposedapproach allows a significant increase of the apparent fracture toughnessvalues.

(a) (b)

(c) (d)

100 mm100 mm

Step of fracture

1 mm

Si3N4 layer

Si3N4–50wt% TiN layer

Crack bifurcation

7.15 Fracture surface of Si3N4/Si3N4–50%wt%TiN composite: (a) and(c) SEI image; (b) and (d) backscattered image.

1 mm


Compressive layerB4C + 30wt% SiC

Tensile layerB4C

B4C + 30wt% SiC

7.16 Schematic presentation of symmetric three-layered design ofB4C/B4C30wt%SiC laminate.

7.5.2 Preliminary ballistic test

The manufactured 90 mm ¥ 90 mm ¥ 10 mm three-layered B4C/B4C–30wt%SiC tiles were tested as armor [67]. The photographs of the experimentset-up of the ballistic test as well as a residual impression in the clay box thatwas used as one of the criteria in the ballistic performance of laminates areshown in Fig. 7.17. The ballistic penetration tests were performed to evaluatethe ballistic performance of the laminates. Depth of penetration tests wereused to evaluate the ballistic performance of the composite laminates. Inaddition, pure B4C monolithic ceramics were used as a standard for the test.Test panels were made using the three-layered B4C/B4C–SiC laminate andB4C monolithic ceramic material as the hard face. While the B4C monolithictile had 100% of its theoretical density, the three-layered B4C/B4C–30wt%SiClaminates had about 3–4% of porosity. A commonly used Spectra fiber-reinforced polymer composite was used as backing plates. The targets weremounted on clay and the projectile was shot at the target at a specific velocity.

Table 7.4 Properties of ceramics used in the stress calculation

Composition E (GPa) Poisson’s ratio CTE (10–6 K–1)

B4C 483 0.17 5.5SiC 411 0.16 3

Table 7.5 Three-layered composite design; total thickness of a tile10.5 mm

Thickness of layers (mm) scomp(MPa) stens(MPa) Apparent

B4C–30wt%SiC B4C KIC (MPa.m1/2)

900 8700 632 131 44


The design of the test panels was selected to ensure defeat of the threat. Thedepth of penetration of the projectile into the backing was measured bypeeling of the unpenetrated layers of the backing plate, and the diameter ofthe impression on the clay after the projectile had been shot was used toevaluate the ballistic performance of the laminate composites. The results ofthe ballistic performance evaluation are shown in Fig. 7.18. As one can seethere was no significant difference in penetration of the projectile into pureB4C monolith ceramics and three-layered composite.

7.5.3 Undesirable influence of tensile residualstresses on a laminate

During the assembly of one 100 mm ¥ 100 mm ¥ 12 mm multilayered tile,the inner thin B4C–30wt%SiC layers were mistakenly replaced with pureB4C thin layers [57]. As a result, instead of a multilayered tile, a three-layered laminate was produced. The parameters of this three-layered tile,

(a)

(b)

7.17 (a) A ballistic test setup; (b) diameter of the residual impressionin the clay box used for evaluation of the performance of the B4Cbased tiles.


including thickness of layers and calculated stresses, are presented in Table7.6. The outer B4C–30wt%SiC layers had a thickness of 1650 mm, and thethick B4C layer had a thickness of 9000 mm. For such a design the level ofresidual tensile stress was raised to 210 MPa after cooling from THP = 2200oC.Such a high residual tensile stress leads to complete fracture of the tileduring decompression of the graphite die to separate the tile after hot pressing(Fig. 7.19). The failure apparently started from the tile edges with crackspropagated further into the tile body.

B4C monolith3-layered composite

1.02 1.04 1.06 1.08 1.1 1.12 1.14Normalized projectile velocity

(a)

0.8

0.7

0.6

0.5

0.4

0.3

0.2

0.1

0

Frac

tio

n b

acki

ng

pen

etra

ted

1.02 1.04 1.06 1.08 1.1 1.12 1.14Normalized projectile velocity

(b)

48

47

46

45

44

43

42

41

40

Cla

y d

efo

rmat

ion

(m

m)

7.18 Ballistic performance results: (a) fraction backing penetrated;(b) clay deformation.

Table 7.6 Three-layered composite design; a total thickness of a tile 12.3 mm

Composition Thickness of layers (mm) scomp (MPa) stens (MPa)

B4C–30wt%SiC B4C

B4C–30wt%SiC/B4C 1650 9000 573 210


This example shows how important it is to determine a critical value ofthe tensile stress in a layer. Certain difficulties exist in finding this criticalvalue. One of the problems is that the mechanical properties of an individuallayer can significantly deviate from those of a corresponding bulk material.We can easily calculate the critical tensile stress if the intrinsic fracturetoughness of a layer and the size of the critical flaw inside the layer areknown, but the critical defect in the layer usually cannot be identified. It ispossible to determine the stress for crack tunneling in the tensile layer [68].Such stress depends only on the intrinsic fracture toughness and the layerthickness. Such transverse cracking of a tensile layer is not possible if thetensile residual stress has a lower value than the stress for crack tunneling.Therefore, an empirical value is used as a critical tensile stress. Such anapproach, in fact, is also rather successful in eliminating cracking in laminates.

7.5.4 Microstructures of three-layered B4C/B4C–SiClaminates

The microstructure of a pure B4C layer of three-layered B4C/B4C–30wt%SiClaminate with 4% porosity is presented in Fig. 7.20. The three-layered B4C/B4C–30wt%SiC tiles tested as armor material had the same microstructureand porosity level as the material shown in Fig. 7.20. As one can see, theporosity at the grain boundary of the ceramics might be a reason why three-layered laminates have not outperformed the dense monolithic boron carbidetiles. A different set of ballistic experiments are required in which fully

7.19 Photograph of fractured three-layered B4C–30wt%SiC/B4C tilehot pressed at 2200oC.


dense boron carbide based laminates will be used for comparison. Suchexperiments will be performed in the future.

A fracture surface of a three-layer tile hot pressed at 2200oC for 1 hourand subsequently broken after hot pressing (shown in Fig. 7.19) is shown inFig. 7.21. The layered composite demonstrates typical brittle fracture. Theinterface between the B4C–30wt%SiC outer layer and the pure B4C innerlayer is shown in Fig. 7.21(a). The fracture surface of the B4C layer ispresented in Fig. 7.21(b). Figure 7.21(c) shows the fracture surface of theB4C–30wt%SiC layer. The cleavage steps on the B4C fracture surface arepresented in Fig. 7.21(d). As one can see from Fig. 7.21, the B4C grain sizein B4C–30wt%SiC layers was in the range of 4–6 mm and the SiC grain sizewas in the range of 2–5 mm. The B4C grain size in pure B4C layers could notbe calculated because of a pure transgranular fracture mode with no grainsor grain boundaries revealed after fracture. Significant grain growth of boroncarbide is expected during hot pressing at 2200oC. However, in B4C–30wt%SiClayers, the existence of the SiC phase prevented the exaggerated grain growthand the grain size distribution was homogeneous. Tiles hot pressed at 2200oCfor 1 hour were fully dense. Tiles hot pressed at 2150oC for 30 or 45 minutescontained some amount of porosity (2–5%) that was concentrated along theinterfaces and mostly in pure B4C layers. Such porosity could be detrimentalto material hardness, affecting Young’s modulus and density, thus significantlylowering the ballistic performance of the laminates. As a result of the hardnessand Young’s modulus decrease, material with a residual porosity of morethan 2% cannot be considered as a candidate for ballistic protection.

While no three-layered composite material was recoverable after the

7.20 Microstructure of pure B4C layer in three-layered B4C/B4C–30wt%B4C composite.


penetration tests shown in Fig. 7.18, in a separate ballistic test designedspecifically at low projectile velocity the debris of three-layered comminutedcomposite was collected to study the microstructural changes in materialafter ballistic impact. The size of the comminuted particles collected afterimpact varied from 5–10 mm to 1–2 mm. The density of this tile was veryclose to the theoretical density of the material, therefore we could considerthat the tile was almost fully dense. The microstructure of the pure B4C layercomminuted by ballistic impact is shown in Fig. 7.22(a). The smooth, flatsurface with transgranular fracture was typically observed for B4C layerswith some amount of cleavage steps present in the material. Such cleavagemode plays an important role both in fracture and in the fragmentation eventduring ballistic impact [69]. There was always some amount of closed porositywhich could not be eliminated by any special treatment such as increase ofhot pressing temperature, pressure or dwell time. The microstructure of theB4C–30wt%SiC layer after impact is shown in Fig. 7.22(b). The B4C grainsare still fractured almost always transgranularly, and a small amount ofclosed porosity was present in boron carbide grains. The fracture surfaces ofall the SiC grains were heavily cleaved, with almost no grains observedwithout cleavage. Such ability of the material to form shear or cleavage steps

(a) (b)

(c) (d)

7.21 Fracture surface of a three-layered tile: (a) interface betweenB4C–30wt%SiC outer layer and pure B4C inner layer; (b) a fracturesurface of B4C layer; (c) fracture surface of B4C–30wt%SiC;(d) cleavage steps on the B4C fracture surface.


should significantly increase the resistance to penetration of SiC ceramiccomposites. This is a topic for further intensive research; however, what isclear at the moment is that both B4C and SiC have distinctively differentdeformation modes under ballistic impact.

This research [57, 64, 67] represents a first step in laminate ceramicsdevelopment that should provide superior ballistic protection. Boron carbide–silicon carbide ceramics have been used in the design and manufacturing ofthree-layered composites with strong interfaces for enhanced fracture toughness.The model of a heterogeneous layered system was used to develop optimaldesign parameters. As a result, laminates with calculated high compressiveresidual stresses (up to 650 MPa) and low tensile residual stresses (below150 MPa) were developed. The feasibility of manufacturing laminate compositesystems with enhanced toughness by incorporation of thin layers with highcompressive stresses in the ceramics was demonstrated. The results of thisstudy are likely to find practical applications in the field of ballistic protectionand mechanical behavior of advanced ceramic composites.

(a)

(b)

7.22 Micrograph of a B4C grain in (a) the pure boron carbide layer,and (b) the B4C–30wt%SiC layer, of the three-layered laminate.Closed porosity was present. Almost all SiC grains have been heavilycleaved.


7.6 Future trends

The most promising approach is the use of layered materials to controlcracks by deflection, microcracking, or internal stresses. In order for thesematerials to become even more useful, their toughness must be increased insuch a way that they could tolerate large flaws during loading. The materialmust be protected against the effect of the largest flaws. The most promisingrecent laminate designs with increased fracture toughness and high residualcompressive stress have been developed with co-workers in the FP5 project‘LAMINATES’ [70]. One of the designs for a B4C/B4C–20wt%SiC laminatebased on these developments is shown in Fig. 7.23. The calculated apparentfracture toughness vs. crack length is also shown in Fig. 7.23. In the proposeddesign the layers create three effective barriers for crack propagation. Theincremental increase of the apparent K1C is specifically targeted in the proposeddesign. With an increasing load the material resistance will grow further andfurther as the crack propagates and, therefore, more energy is required forlaminate fracture. The apparent fracture toughness for this design has a

7.23 Design of B4C/B4C–20wt%SiC for increased fracture toughness.

Crack

B4C

B4C

B4CB4C + 20wt% SiC

0 1 2 3 4Crack length (mm)

14

12

10

8

6

4

2

0

Ap

par

ent

frac

ture

to

ug

hn

ess

(MP

a m

1/2 )


maximum calculated value of ~14 MPa m1/2 while the intrinsic fracturetoughness of B4C was adopted to be only 2 MPa m1/2. The thick inner B4Clayer serves to obtain a low level of tensile residual stress.

The following directions of research in mechanical behavior improvementof ceramic laminates are proposed:

∑ Improvement of the material’s structure in separate layers∑ Obtaining macrocrack shielding∑ Obtaining macrocrack deflection∑ Development of gradient laminar structures.

Opportunities for structure optimization of layered composites are:∑ Optimization of statistical parameters of the strength distribution of

composites∑ Optimization of microcracking process in individual layers∑ Design of layered structures with maximum apparent fracture toughness

induced by crack shielding∑ Design of layered composites with crack bifurcation in compressed layers∑ Control of residual stress distribution in laminates∑ Design of asymmetric laminar structures tailored for specific use.


This work was supported by the European Commission, project 1CA2-CT-2000-10020 Copernicus-2 ‘Silicon nitride based laminar and functionallygradient ceramics for engineering application’, and by AFOSR, project F49620-02-0340. EMPA was funded by BBW, the Swiss Federal Office for Educationand Science, under contract 99.0785. This work was also partly performed atthe Army Center for Nanoscience and Nanomaterials, North Carolina A&TState University.

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J. Am. Ceram. Soc., 62(3–4), 149–153, 1979.51. Niihara, K., Nakahira, A., Hirai, T., The effect of stoichiometry on mechanical

properties of boron carbide, J. Am. Ceram. Soc., 67, C13, 1984.52. Thevenot, F., in Properties of Ceramics, ed. G. de With, Terpstra, R.A., Metselaar,

R., Elsevier Applied Science, London and New York, 1989.53. Lee, H., Speyer, R., Hardness and fracture toughness of pressureless sintered boron

carbide (B4C), J. Am. Ceram. Soc., 85(5), 1291–1293, 2002.54. Wilkens, M.L., Mechanics of penetration and perforation, Int. J. Eng. Sci., 16, 793–

807, 1978.55. Hyatt, T.P., Electronics: tape casting and roll compaction, Am. Ceram. Soc. Bull.,

74(10), 56–59, 1995.56. Dec, R.T., Zavaliangos, A., Cunningham, J.C., Comparison of various modelling

methods for analysis of powder compaction in roller press, Powder Technology, 130,265–271, 2003.

57. Orlovskaya, N., Kalidindi, S., Lugovy, M., Subbotin, V., Radchenko, O., Adams, J.,Chheda, M., Shih, J., Sankar, J., Yarmolenko, S., Robust design and manufacturingof ceramic laminates with controlled thermal residual stresses for enhanced toughness,J. Mat. Sci. (in press).

58. Kuebler, J., Fracture toughness using the SEVNB method: preliminary results, Ceram.Eng. Sci. Proc., 18, 155–162, 1997.

59. Kuebler, J., Fracture toughness of ceramics using the SEVNB method: from apreliminary study to a standard test method, in Fracture Resistance Testing of Monolithicand Composite Brittle Materials, ASTM STP 1409, ed. J.A. Salem, M.G. Jenkinsand G.D. Quinn, ASTM, West Conshohocken, PA, pp. 93–106, January 2002.

60. Hermann, M., Shubert, C., Pabst, J., Richter, H.J., Obenau, P., Yaroshenko, V. inSymposium: Verstaekung Keramisher Werkstoffe, Hamburg, October 1991,Tangungsband, DGM.

61. Orphal, D.L., Franzen, R.R., Charters, A.C., Menna, T.L., Piekutowski, A.J., Penetrationof confined boron carbide targets by tungsten long rods at impact velocities from 1.5to 5.0 km/s, Int. J. Impact Eng., 19, 15–29, 1997.

62. Shih, C.J., Meyers, M.A., Nesterenko, V.F., Chen, S.J., Damage evolution in dynamicdeformation of silicon carbide, Acta. Mater., 48, 2399–2420, 2000.

63. Orphal, D.L., Franzen, R.R., Penetration of confined silicon carbide targets by tungstenlong rods at impact velocities from 1.5 to 4.6 km/s, Int. J. Impact Eng., 19, 1–13,1997.

64. Orlovskaya, N., Lugovy, M., Subbotin, V., Radchenko, O., Adams, J., Chheda, M.,Shih, J., Sankar, J., Yarmolenko, S., Design and manufacturing B4C–SiC layeredceramics for armor applications, in Ceramic Armor and Armor Systems (CeramicTransactions, Volume 151), 2003, pp. 59–70.


65. Kislyi, P.S., Kuzenkova, M.A., Bondaruk, N.I., Grabchuk, B.L., Boron Carbide,Naukova Dumka, Kiev, 1988 (in Russian).

66. Orlovskaya, N., Lugovy, M., Kuebler, J., Mechanical performance of 3 layeredB4C–SiC ceramic composites, unpublished results.

67. Orlovskaya, N., Adams, J., Chheda, M., Shih, J., Yarmolenko, S., Sankar, J., Lugovy,M., Subbotin, V. Boron carbide–silicon carbide laminate ceramics for ballisticprotection, Proc. of 2003 ASME Int. Mech. Eng. Cong., Vol. 3, Paper IMECE2003-43323, 15–21 November 2003, Washington, DC.

68. Ho, S., Suo, Z., Tunnelling cracks in constrained layers, J. Appl. Mech., 60, 890–894, 1993.

69. Chen, M., McCauley, J.W., Hemker, K.J., Shock-induced localized amorphization inboron carbide, Science, 299, 1563–1566, 2003.

70. Final Report 2004, EC/BBW Contract No. ICA-CT-2000-10020, FP5 INCO-Copernicusproject ‘LAMINATES’ (Silicon Nitride Based Laminar and Functionally GradedCeramic Composites for Engineering Applications), project partners: University ofWarwick (UK), FCT Technologie (Germany), Institute for Problems of MaterialsScience (Ukraine), Materials Research Center Ltd (Ukraine), Institute for Problemsof Strength (Ukraine), Institute of Chemical Physics (Armenia), Drexel University(USA), EMPA (Switzerland).

216

8.1 Introduction

There is a growing demand for ultra-hard materials in applications such ashigh-speed cutting and forming, hard disc drives and various optical andbiomechanical components. Normally properties other than hardness arealso needed, such as chemical inertness, a low friction coefficient and athermal expansivity with respect to that of the substrate so the coating doesnot spontaneously peel off. The resulting compromise is such that the transitionmetal nitrides, with hardnesses of approximately 20 GPa are often used,rather than the very hardest materials such as diamond (100 GPa) or cubicboron nitride (50 GPa). There is therefore considerable interest in developingways of increasing the hardness of an intrinsically hard material by modifyingits microstructure. The problem is not a trivial one. Most of the strengtheningmechanisms that have been developed in metals give rise to an increment ofstrength that is independent of the strength of the base alloy. The total strengthis simply the sum of the initial strength and the increment, which has atypical magnitude of a few hundred megapascals. However, in materials thatare intrinsically strong, such as most ceramics, the resistance of the latticealone to dislocation motion gives materials with strengths of tens of gigapascals.The strengthening mechanisms used in metals are therefore normally quiteuseless.

In the late 1980s it was shown that a multilayer made of alternating layersof TiN and VN, each layer a few nanometres in thickness, could showhardnesses of over 50 GPa, twice that of the monolithic material (Helmerssonet al., 1987). Since then a number of ideas have been developed to explainthe strengthening that is observed, such as coherency strains or changes indislocation line energy between the layers. However, there are such structuresthat on the basis of these analyses might be expected to show hardening butdo not do so, suggesting either that some other effect is leading to weakeningor that the hardening effect does not arise as has been suggested. The aim ofthis chapter, therefore, is to review what is known about how multilayer

8Hardness of multilayered ceramics

W J C L E G G, F G I U L I A N I, Y L O N G,S J L L O Y D, University of Cambridge UK and

J M M O L I N A - A L D A R E G U I A, Centro de Estudiose Investigaciones Tecnicas de Gipuzkoa (CEIT), Spain

Hardness of multilayered ceramics 217

ceramic coatings deform, why they might be hard and why this hardening isnot reliably observed.

8.2 Behaviour of multilayer structures

The brittleness of the ceramic coatings has meant that their flow behaviourhas been inferred from hardness tests rather than directly from tensile tests.Results from a large number of different systems are shown in Figs 8.1(a)–(c), where it can be seen that a very wide range of properties are found. Insome, large increases in hardness are found (Helmersson et al., 1987; Mirkarimiet al., 1990; Shinn et al., 1992), well above those that might be expectedfrom the properties of the individual components. In others, there is no effect(Ljungcrantz et al., 1998; Yashar and Barnett, 1999; Högberg et al., 2001;Molina-Aldareguia et al., 2002; Barnett et al., 2003). Even reductions inhardness have been observed (Jayaweera et al., 2003). Rather surprisinglythe greatest increases in hardness have been observed in systems where thephases are isostructural, such as TiN and NbN, both of which have the B1rocksalt structure: see Fig. 8.1.

0 10 20 30 40Wavelength (nm)

(a)

60

50

40

30

20

10

0

Har

dn

ess

(GP

a)

NbNVC

NbN, LNbN,B

TiC

TiN

TiN-VN (Helmerssonet al., 1987)

TiN-NbN (Shinn et al.,1992)

TiN-VNbN (Mirkarimiet al., 1990)

NbN-VNbN (Shinn &Barnett, 1994)

TiN-NbN B-series(Ljungcrantz et al., 1998)TiN-NbN L-series(Ljungcrantz et al., 1998)

TiN-NbN (Molina-Aldaregnia et al., 2002)

TiC-VC (Högberg et al.,2001)

8.1 The variation of hardness with multilayer wavelength in a rangeof different types of structures. These include multilayers of (a)isostructural transition metal nitrides and carbides, which show thegreatest hardening; (b) nonisostructural multilayer materials, whereslip cannot occur by the movement of dislocations across the planesof the composition modulation, because the slip systems aredifferent in the two materials; and (c) materials where differentcrystal structures are stabilized at small layer thicknesses, such asAlN deposited onto TiN.


The use of indentation complicates the interpretation of the measurements.However, relationships between hardness and flow stress exist for monolithicmaterials and these have been used to obtain information in the multilayers.When an indenter is pressed into the surface of a material, the material thatis displaced must be accommodated either by material being pushed out of

Cubic AlN fully stabilized by TiN

Setoyama et al., 1996

Wong et al., 2000Li et al., 2004

0 10 20 30 40 50Wavelength (nm)

(c)

40

30

20

10

0

Har

dn

ess

(GP

a)

Mo-NbN (Madan et al., 1998)Y2O3-ZrO2 (Yashar & Barnett 1999)TiN-TiB2 (Barnett et al., 2003)


(b)

60

50

40

30

20

10

0

Har

dn

ess

(GP

a)

TiB2

TiN

NbNYSZ

Mo

Y2O3

8.1 Continued


the surface (pile-up) (Tabor, 1951) or by moving radially outward from theindentation so that it is accommodated elastically within the body (Marsh,1963). The former tends to occur in metals whilst the latter occurs in materialswith a high ratio of the uniaxial flow stress, sf, to the Young modulus, E.These include most ceramics, which are the subject of this review. Numericaland analytical solutions indicate that for a soft metal the hardness, H, shouldbe about three times the uniaxial flow stress, sf, falling gradually to about 1as the ratio of H/E increases (Marsh, 1963; Johnson, 1970; Cheng and Cheng,2000).

This radial flow requires the movement of atoms across the planes of thecomposition modulation. Where the two materials have very similar crystalstructures, as for instance in strained layer superlattices, where the interfacesbetween the layers are coherent, dislocations may be able to simply glideacross the layers. In this case the dislocation must move under conditionswhere the coherency or elastic misfit strains between the layers alternate,giving rise to stresses that act on the dislocation. However, the changes inelastic modulus and Burgers’ vector cause changes in the dislocation lineenergy as the dislocation moves and hence to a force on the dislocationarising from the composition modulation, giving a possible source of hardening.However, in the more general case where the layers are not isostructural andso do not have common slip systems, slip will have to be renucleated in eachlayer as envisaged by Hall and Petch, even though the layers may be coherent.Alternatively, where flow across the layers becomes difficult, then thedeformation required to accommodate the indentation might occur by thelateral movement of material within the individual layers. To investigate thebehaviour of ceramic multilayers we therefore need to examine the variousmechanisms and understand under what conditions of structure and appliedindenter pressure such a mechanism might provide the dominant obstacle todeformation.

8.3 Hardening mechanisms in multilayers

8.3.1 Hardening due to coherency stresses

The initial observations of hardening in ceramic multilayers were made instrained layer superlattices of TiN/VN (Helmersson et al., 1987). At multilayerwavelengths of approximately 10 nm, the hardness was greater than 50 GPa,much greater than the 20–25 GPa normally reported for thin films of monolithicTiN. Explanations focused on the interaction of the elastic stress field of thedislocation with the alternating stress field of the superlattice, which givesrise to shear stresses that are a maximum on planes running through thethickness of the layers in both in-plane directions at an angle of 45∞ (Cahn,1963). These shear stresses alternate so that one layer will aid the passage of


a given dislocation, whilst the next will impede it. The effect of such alternatingstress fields on the movement of a single dislocation has been analysed byCahn (1963) and later modified by Kato et al. (1980). They analysed thesituation of three orthogonal waves of a sinusoidal composition modulation,as might occur in a face centred cubic alloy undergoing spinodal decomposition.They found that the increment of the shear flow stress, Dtc, due to thevarying internal stress field could be given approximately by

Dtc = 0.14AEh (8.1)

where A is the magnitude of the composition amplitude, E is the Youngmodulus and h is the misfit strain, assuming that the lattice parameter varieslinearly with changing composition. Setting A to 1, that is the compositionchanges from pure TiN to pure VN, taking values of the lattice parameter ofTiN and VN to be respectively 0.424 nm and 0.414 nm (see Table 8.1) andusing the Young modulus of TiN as 450 GPa, gives an increment in the shearflow stress of approximately 1.5 GPa, suggesting an increase in the hardnessof about 9 GPa, much smaller than the increment observed in Fig. 8.1(a).However, the analysis is for a composition modulation in three orthogonaldirections rather than the one direction here.

The simple geometry in the multilayers gives rise to a further complication.For a multilayer made of two materials with different lattice parameters, theatom spacings parallel to the planes of composition modulation will be equalif the layers are coherent. However, normal to the planes, the atom spacingswill be unchanged, apart from a Poisson contraction that will act to increasethe difference, so that the Burgers’ vectors in the two materials will mostlikely be different. Kelly (1991) has pointed out that when a dislocationmoves from one phase, B, to another, A, a dislocation with a Burgers’ vectorgiven by the difference at the Burgers’ vectors between the two phases willbe left behind in the interface. This will give rise to a repulsive force on theglide dislocation when it is moving in the layer with the smaller Burgers’vector, as the two dislocations will have the same sign. However, when theglide dislocation is moving through the layer with the larger Burgers’ vectorand the interface dislocation has an opposite sign, the force between the twodislocations will be attractive, inhibiting the movement of the glide dislocation.The force between the two dislocations can be estimated using the standard

Table 8.1 Single crystal elastic constants and lattice parametersof TiN, NbN and VN. The data are taken from Kim et al. (1992)

Material c11 (GPa) c12 (GPa) c44 (GPa) a (nm)

TiN 625 165 163 0.424NbN 556 152 125 0.439VN 533 135 133 0.414


expressions from dislocation theory; thus it is predicted to increase as thedistance between the dislocations, r, is diminished, and to rise without limitas r tends to zero. To avoid this difficulty, Kelly follows Koehler (1970) andassumes that the force will be at a maximum when r = 2b, where b is theBurgers’ vector. Substituting for r gives the maximum force on the dislocation,and as F = tb, the increase in the shear flow stress, Dtm, required to move thedislocation a long distance through the superlattice against the forces due tothese misfit dislocations is given by

Dt pD

mA

A =

2

G bb

◊ (8.2)

where GA is the shear modulus of layer A, bA its Burgers’ vector and Db thedifference between the Burgers’ vectors. Using the values of the latticeparameters of TiN and VN and the shear modulus of TiN, as before, gives anincrease in the shear flow stress of 0.9 GPa, and a corresponding incrementin the hardness of approximately 4.5 GPa, which again is much less than thehardening observed: see Fig. 8.1(a).

One difficulty in interpreting the results in terms of coherency strains isthat as the layer thickness increases, coherency is lost and an array of misfitdislocations is formed. This phenomenon has been studied in great detail insemiconductors, where retaining coherency is of considerable importance.Dunstan and co-workers have studied the effects of coherency stresses usingstrained layer superlattices of InGaAs where the differences in lattice parameterbetween the layers are introduced by doping layers with varying amounts ofIn (Jayaweera et al., 2003). Looking at layers with thicknesses ranging from17 to 125 nm and strain modulations varying from 0 to 1.46%, they foundthat the introduction of coherency stresses weakened rather than strengthenedthe material. Furthermore it appeared that the magnitude of the strengthreduction was not dependent on either the strain modulation or the extrastrain energy associated with the strained layers. Rather the hardness wasgiven by an expression of the form

H H kMFY YO = – ¢ (8.3)

where HY is the measured hardness of the superlattice, HYO is the hardness

of the monolithic material, k is a numerical constant, M is a biaxial elasticmodulus and F¢ is a thickness-averaged strain modulation given by theexpression

¢Fh hh h

= | + |

+ c c t t

c t

e e(8.4)

where ec and et are the strains in the compressive and tensile layers respectivelyof thicknesses hc and ht. By considering the rate of change of elastic energywith respect to an unspecified variable, it is shown how an expression of this


form might be obtained. The formulation in terms of a strain energy gradientinvolving both layers is taken to imply that the flow process must requiresome minimum volume to operate, although it is not clear why such effectsdo not also appear in the other experiments, or even in spinodal structureswhere the microstructural features are even smaller.

It is tentatively suggested that this unknown variable may be the rate ofplastic work, although it is not clear how a yield criterion can be developedby considering the rate of change of a variable with respect to the rate ofplastic work. Despite this, as the authors point out, the formal derivation iscorrect regardless of the physical meaning of the unknown variable. However,the derivation does predict that the sign of F¢ is important, despite equation(8.4). Unfortunately there are no experimental results that might resolve this.

In summary it is clear that the effects of coherency stresses are very farfrom understood, although experiments suggest that in intrinsically strongmaterials their effect is relatively small. This has led to the consideration ofother possibilities, in particular the effects of elastic inhomogeneity.

8.3.2 Hardening due to changes in dislocationline energy

The importance of such effects was deduced from the data shown in Fig. 8.2from the elegant experiments of Barnett and co-workers (Chu and Barnett,1995). In the TiN/VN multilayers the two layers are both elastically strainedwith respect to one another and have different elastic moduli, that of TiNbeing greater than that of VN, which is similar to that of NbN. Shinn andBarnett (1994) have used this to study the effects of elastic modulus mismatch.As shown in Fig. 8.2, systems where there was a difference in elastic modulusshowed a substantial increase in hardness. Where there was no difference inelastic modulus little or no hardening was observed, whilst hardening wasobtained in a TiN/V0.6Nb0.4N system where there was a modulus mismatchbut no lattice mismatch (Mirkarimi et al., 1990; Hubbard et al., 1992).

These observations have a theoretical foundation in Cahn’s original analysisof hardening in spinodal structures (Cahn, 1963). A dislocation has a lineenergy, U, associated with its elastic misfit in the lattice whose magnitude isapproximately equal to the product of the shear modulus, G, and the squareof the Burgers’ vector. If the body containing the dislocation is not uniformbut is instead made up of layers of different materials whose values of G andb give different values of U, there will be a force on the dislocation acting tomove it from the layer with the higher value of U to that with the lowervalue.

The situation of a screw dislocation moving towards an atomically sharpinterface separating two isotropic layers, A and B, with shear moduli GA andGB respectively but the same Burgers’ vector, was analysed by Koehler


(1970) using the image force approach (Benlahsen et al., 1993), therebyimplicitly assuming that the Burgers’ vectors are the same in each material,or at least that the effects of differences in Burgers’ vector are unimportant.Assuming linear elasticity, Koehler showed that the magnitude of this forcedepended on the distance, r, between the dislocation and the interface betweenthe two layers and the elastic mismatch across the interface, Q, given by

QG GG G

= – +

A B

A B(8.5)

As might be expected, the magnitude of the force on the dislocation increasesas the dislocation gets closer to the interface. The force is repulsive when thedislocation is moving towards the stiffer layer, and attractive when it ismoving towards the more compliant layer.

The major difficulty with the Koehler solution is that the force on thedislocation is predicted to rise without limit as the dislocation reaches the

TiN-VN (Helmersson et al., 1987)

TiN-NbN (Shinn et al., 1992)TiN-VNbN (Mirkarimi et al., 1990)NbN-VNbN (Shin and Barnett, 1994)


Loop motion

Sourceactivation

60

50

40

30

20

10

0

Har

dn

ess

(GP

a)

8.2 The variation in hardness theoretically predicted for the TiN/NbNmultilayers compared with data for the isostructural nitrides shownin Fig. 8.1(a). The predictions are based on the ideas of an incrementof hardness arising from the elastic mismatch across the layers,equation (8.6), shown as the horizontal dashed line, and for thelateral flow of material within the interlayers, equations (8.10–8.12),shown as the dashed line increasing as the wavelength decreases. Inthe isostructural multilayers, the upper limit to the increase inhardness should occur when the stress is high enough to drivedislocation motion across the layers. In the nonisostructural, thiscondition does not apply, or is substantially modified. The linesshown assume that the hardness of monolithic TiN and NbN is 25GPa and that the layer thicknesses are equal. Other data are given inTable 8.1.


interface and r = 0. To obtain a solution, Koehler assumed that the magnitudeof the force on the dislocation due to the interface would reach a maximumwhen r = 2b. However, the magnitude of the force has a maximum valuewhen, for a given increment of movement towards an interface, there is thegreatest change in the elastic strain energy due to the dislocation misfit, thatis when the dislocation is crossing the interface and r = 0.

A solution to this problem requires a knowledge of the atom positions inthe dislocation core. Using the atom positions in the Peierls dislocation,Pacheco and Mura (1969) estimated the force on a dislocation due to just asingle interface. They obtained the increase in shear flow stress, DtE, due toan elastic modulus change across a sharp interface as

Dtp

qE

B2 = 2 sin

GQ (8.6)

where q is the angle of the slip plane with the interface and the dislocationis in material B and is greater than that predicted by Koehler by a factor of16p. An estimate of the magnitude of this effect was obtained by Shinn andBarnett, setting the shear modulus equal to the single crystal elastic constantc44 (Shinn and Barnett, 1994; Chu and Barnett, 1995). These are 163 and 125GPa for TiN and NbN respectively: see Table 8.1. This gives a value of Q of0.14 and a value of DtE of 5.3 GPa. Taking the hardness, H, to be three timesthe uniaxial flow stress and this to be twice the shear flow stress, and followingChu and Barnett (1995), the increment in the hardness, DH, due to the elasticmismatch of the layers is predicted to be 48 GPa. This assumes that thecomposition varies from pure TiN to pure NbN. However, X-ray diffractionshowed that the variation was only one-half of this (Shinn and Barnett,1994), so that the predicted value of DH was approximately 24 GPa and invery reasonable agreement with the data.

However, in a cubic structure the value of G will be equal to c44 onlywhen slip is on the {110}<001> slip system (Kelly et al., 2000). In rocksalt-structured nitrides and carbides, slip in indentation at room temperatureoccurs on the {110} <110> slip system (Williams and Schaal, 1962; Molina-Aldareguia et al., 2002). The appropriate value of G is related to the differentsingle crystal elastic constants, cij, by

G c c = 12

( – )11 12 (8.7)

Substituting the relevant values for TiN and NbN from Table 8.1 gives avalue for Q of only 0.06, one-half of that obtained above. Accounting for theobserved concentration modulation, a total increment in hardness of only 12GPa is predicted. If the uniaxial flow stress is taken as twice the shear flowstress, the predicted value of DH falls to only 9 GPa, much less than thatobserved here.


At higher temperatures flow occurs on the {111} <110> slip system(Williams and Schaal, 1962), where G is given by (Kelly et al., 2000)

Gc c cc c c

= 3 ( – )4 + –

44 11 12

44 11 12(8.8)

However, even here Q is only 0.09, giving a DH of approximately 14 GPa,rather than the 24 GPa given by Shinn and Barnett (1994).

Only one interface has been considered. In the multilayer, there will beforces due to each of the interfaces, each decreasing as one moves furtheraway from the dislocation. When the layers are thick this effect is negligible,so the increase in stress required to overcome the elastic mismatch across theinterface will be given by equation (8.6). However, when the layers are verythin, the force due to the elastic mismatch will be reduced. As an approximation,Lehoczky considered the force on the dislocation due to just the two interfaceslying on either side of the dislocation (Lehoczky, 1978b). Adding moreinterfaces decreases this but, as they are further away than the two nearestinterfaces to the dislocation, their effect is small and a solution includingmany interfaces is within 5% of that for just two interfaces. However, theeffect becomes important, at least for the systems here, only when the multilayerwavelength falls below about 5 nm. There are therefore two effects thatcause a reduction in the hardness of superlattices with very fine layers. Thefirst is due to the inevitable intermixing that occurs at the interface, an effectthat becomes more marked as the layers become thinner (Shinn and Barnett,1994). The second is due to the presence of other interfaces. Both would actto reduce the hardness.

These analyses describe the situation where flow occurs by the motion ofsingle dislocations passing across the interfaces between the two materials.However, if the layers are very thick, dislocation sources within the individuallayers will be able to operate. The shear stress required to operate a sourcein the more compliant layer, B, is given by

tSB B

B ª G b

l(8.9)

where lB is the dimension of the source. If this stress is less than the repulsiveforce acting on it due to the presence of the high modulus layer, then thedislocations will pile up at the interface, giving rise to a stress concentrationat the head of the pile-up, tpu, which will help force the dislocation throughthe interface (Lehoczky, 1978a). The flow stress of the layer will thereforeincrease as the layer decreases in thickness until the flow stress has increasedby DtE, at which point pile-ups will be unable to form and the increase inflow stress will be that given by equation (8.6). This is shown in Fig. 8.2 fora TiN/NbN multilayer, where the hardnesses of monolithic TiN and NbN areboth taken as 25 GPa and the layer thicknesses of each are assumed to beequal.


None of this explains why hardening is observed in one set of TiN/NbNdata (Shinn et al., 1992) but not in three others (Ljungcrantz et al., 1998;Molina-Aldareguia et al., 2002): see Fig. 8.2. Nor does it explain the lack ofhardening in the TiC/VC system (Högberg et al., 2001). All of these have thesame crystal structure and were single crystal, with the exception of the dataof Ljungcrantz et al., where the layers were polycrystalline. Ljungcrantz’sdata is interesting because it contains two materials, denoted B and L, madein different laboratories. Apart from a slight difference in hardness betweenthe two, neither showed any hardening.

One possibility is differences in the sharpness of the interfaces. The forceon the dislocation due to the elastic inhomogeneity is dependent upon therate of change of the elastic constants and hence on the rate of compositionchange. Ideally this should be abrupt, but it rarely is. The magnitude of thiseffect has been estimated by Krzanowski (1991, 1992). By determining howthe composition, and hence the modulus, changed for a small increment ofmovement, he was able to show that the effect of increasing the interfacialwidth from 1 nm to just 3 nm would decrease the increment of flow stress byapproximately an order of magnitude. Using X-ray diffraction and assumingthat the composition would change linearly across the interface betweenlayers of pure TiN and NbN, Chu and Barnett (1995) measured the interfacewidths to be approximately 2 nm. Using Krzanowski’s analysis they foundthat the hardness increment predicted was no greater than that obtainedpreviously, where it was assumed that the compositional changes were dueentirely to the composition modulation. In any case it does not explain whysometimes no increase in the hardness is observed as the interface widths insome experiments were measured to be approximately 1 nm (Molina-Aldareguia et al., 2002), less than those measured by Chu and Barnett (1995)in their experiments. Sharp X-ray satellite peaks indicative of sharp interfaceswere also measured in the other work (Högberg et al., 2001).

Furthermore it is not clear why no hardening is observed in the NbN/VNmultilayers. Whilst it is true that there is little difference in G for NbN andVN, the difference in dislocation line energy is associated with a change inthe dislocation line energy, U, which is proportional to Gb2, taken as|( – | /( + ).A A

2B B

2A A

2B B

2G b G b G B G b Using Table 8.1 and remembering thatTiN, NbN and VN have the same crystal structures and slip systems, it canbe seen that although the elastic mismatch, Q, is indeed much smaller in thecase of the NbN/VN (0.07%) multilayer than it is for the TiN/NbN (6.5%)multilayer, the difference in the value of Gb2 is reversed, the differencebeing only 6% for TiN/NbN but 13% for NbN/VN. It is therefore not clearthat the absence of any hardening in the NbN/VN system is consistent withthese ideas as claimed.


8.3.3 Hardening due to lateral flow of material

It has also been suggested that flow might occur at lower stresses than thosepredicted above by movement of material within the individual layers (Chuand Barnett, 1995). This has been observed in pearlitic structures made up ofalternating layers of ferrite and cementite, and observations in other multilayersystems suggest that that deformation might occur in this way (Gil-Sevillano,1979). Two cases have been identified: the first where only the movement ofa pre-existing dislocation loop is required, the second where the activation ofa dislocation source within the layer is needed. Gil-Sevillano (1979) showedthat the extra stress, DtM, required to move a dislocation half-loop in a layerof width l is

Dt a qqM =

2 cos ln

cos Gb

ll

b◊ Ê

Ëˆ¯ (8.10)

where a = 1/4p and q is the angle between the slip-plane and the normal tothe interface. The extra stress required to activate a dislocation, DtA, is twicethat for motion. The overall shear flow stress of a given layer is thereforeequal to DtM plus any contribution from the lattice resistance, tL, which inthese materials in bulk form is the dominant contribution to the flow stress.Chu and Barnett (1995) then assume that the multilayer of materials A andB is strained as it were in compression, so that the overall uniaxial flowstress is given by the volume-averaged flow stress, that is

s s s = + AA

BBl l

L L (8.11)

where

s t DtA,B L M = ( + )¢m (8.12)

where m¢, the constant relating the uniaxial to the shear stress, is taken as 2.The overall hardness of the multilayer can then be obtained by taking thehardness as three times the uniaxial flow stress, taking the Burgers’ vector tobe a/2÷2, where a is the lattice parameter, as slip occurs on the {110}<1 1 0>slip system, and taking an average value of cosq to be 0.5.

The predicted hardnesses given by both dislocation motion and sourceactivation are shown in Fig. 8.2. Comparing this with the data on TiN/NbNmultilayers, it can be seen that the rising portion of the curve is a fair fit tothe data of Shinn et al. (1992). However, it is not clear why such a processmight operate at a considerably higher stress than that required to drivedislocations across the interfaces, as discussed in the previous section andshown as the horizontal line in Fig. 8.2. Indeed the predictions seem moreconsistent with the data where little or no hardening was observed.

The possibility that hardening might arise simply because flow is restricted


to the individual layers greatly increases the range of systems that might bedeveloped. The materials considered so far have generally been isostructural(generally all having the B1 rocksalt structure). However flow can be restrictedto the individual layers simply by using materials which have dissimilarstructures so that flow cannot take place simply by the movement of adislocation from one phase to the next. A range of systems have beeninvestigated, including TiN/TiB2, ZrN/ZrB2 (Barnett et al., 2003) and systemscontaining metal layers, such as Mo/NbN, W/NbN (Madan et al., 1998),TiN/Cu (Ljungcrantz, 1995) and Y2O3/ZrO2 (Yashar and Barnett, 1999).Some hardening is seen in the Mo/NbN and W/NbN systems, but the effectis much less pronounced than in the TiN/NbN described earlier and nohardening at all is observed in the other systems: see Fig. 8.1(c). The reasonsfor this are not clear, as in all cases the interfaces are extremely sharp(Barnett et al., 2003).

It can be seen from equations (8.10) and (8.11) that the contribution to thehardening comes mainly from the layers with a higher elastic modulus.However, differences in the elastic properties between the layers will causethe loops in the stiffer layers to be pulled across the interfaces, for the samereason that loops in the less stiff layers are repelled by the interfaces, greatlydiminishing their contribution to the overall flow stress.

Furthermore the assumption that the overall hardness is given by a thickness-averaged hardness of the two materials applies if the multilayer were to bepulled in tension with the layers parallel to the tensile axis, but it is not clearwhether this is true when the sample is being indented. For instance there areobservations that, under a nanoindentation, the deformation is concentratedin the weaker of the two phases, as shown in Fig. 8.3, which shows a cross-section through an indentation in a TiB2/Al multilayer, where the compressivestrain in the Al layers is greater than that in the TiB2 layers. The hardness isobserved to fall very rapidly as the metal volume fraction is increased,although there is still deformation in the TiB2 layers, which may occur dueto porosity in the layers. The wavelength in this material is 200 nm, whichis relatively thick.

In a TiN/Cu multilayer where L = 4.5 nm (Ljungcrantz, 1995), the hardnessvaries between the values for monolithic films of TiN and Cu, suggestingthat the metal layer is indeed constrained when it is very thin, (see Fig. 8.4)but that there is no need to invoke atomistic processes. Only in the Mo/NbNmultilayers (L = 5 nm) does the hardness rise above that even where themetal fraction is high, although the effect is not large (Barnett et al., 2003).The reasons for the difference between the TiN/Cu and the Mo/NbN multilayersare not clear, but are possibly associated with the difficulty of plastic flow inboth molybdenum and tungsten. However, the effect is still not large. And asradial flow is preferred in monolithic materials, one might expect that thelateral flow occurring in these multilayers might be more difficult, giving


500 nm

8.3 Deformation under an indentation in an Al/TiB2 multilayer, with awavelength of 200 nm, showing that deformation is concentrated inthe Al.

8.4 The effect of varying the thickness fraction of the phases inmultilayers of Al/TiB2, TiN/Cu and Mo/NbN, with wavelengths of 200,4.5 and 5 nm respectively. In the coarser multilayer, deformationappears to be concentrated in the softer layer. In the TiN/Cumultilayer, the behaviour is close to a rule of mixtures, whereas forthe Mo/NbN multilayer, some further hardening appears to beoccurring.

0 0.5 1Metal fraction

30

20

10

0

Har

dn

ess

(GP

a)

Cu filmMo film

= Mo/NbN= TiN/Cu= Al/TiB2

TiB2 film

TiN filmNbN film


some hardening consistent with that observed. Unfortunately the magnitudeof such an effect is unknown due to the lack of any theoretical studies.

8.3.4 Summary

It can be seen that the idea of flow occurring within the individual layersmight explain why some materials are harder than either of the materialsfrom which they are made, giving a fair fit to the analyses. However, mostmultilayers show hardnesses in between those of the individual componentsor with only minimal hardening, including some TiN/NbN multilayers, eventhough the interfacial widths are measured to be the same as in those wherehardening is observed, as well as in carbide and in oxide superlattices. Whilstatomistic mechanisms may be required to explain the magnitude of the hardnessin some cases, it appears that in many they do not. Despite this, the observedlateral flow is expected to give rise to some hardening, even where flow istreated in a continuum manner.

8.4 Microstructural changes due to

making a multilayer

So far we have considered the properties of a multilayer only in terms of theeffect of the composition modulation on the movement of dislocations eitheracross the layers of different composition or, separately, within them. Thisgives a certain measure of agreement with particular sets of data, but whenthe experiments are considered in their entirety, these ideas cannot accountby themselves for the observed behaviour. This is noticeable in Fig. 8.1(a),which shows the data for the isostructural nitrides and carbides, where thereappears to be a significant differences in the hardnesses of the multilayers atlonger wavelengths and in the monolithic films, suggesting the importanceeither of internal stresses or of microstructural changes to the individuallayers of the multilayer, caused by making the material in a multilayer form.

8.4.1 Hardening due to internal stresses

It is well known from the work of Thornton, Hoffman and others that largeresidual stresses can be developed in thin films (Thornton and Hoffman,1989; Hoffman, 1994). These may arise due to expansivity mismatch withthe substrate or due to stresses induced during the growth of the coating,which can be varied depending on the conditions of temperature atmosphereand deposition method under which the film is grown. For more informationon the origin of stresses in thin films, the reader is referred to the reviews byWindischmann (1987, 1992). Furthermore, because the coatings need typicallybe only a few microns in thickness, they can sustain internal stresses of the


order of gigapascals without peeling off the substrate (Kendall, 1975), althoughspecial precautions may be taken to improve the adhesion of the coating.

There is considerable experimental work in both metallic and ceramicfilms, showing that such internal stresses can greatly increase the measuredhardness. The simplest explanations are in terms of superimposing an in-plane stress on the overall maximum shear stress under the indenter, althoughPharr (Bolshakov et al., 1996; Tsui et al., 1996), looking at nanoindentationof Al films, considers the effect to arise due to pile-up around the indenter,causing the actual depth of penetration, and hence area of the indentation, tobe greater than that calculated.

However, regardless of the origin, residual compressive stresses in filmsare clearly associated with increases in the measured hardness of bothmonolithic and multilayer films. Figure 8.5 shows the results of Münz andhis co-workers for different multilayers but all with a bilayer period ofbetween 3 and 4 nm (Lewis et al., 1999; Münz et al., 2001). It can be seenthat hardnesses of up to 70 GPa have been measured. Also plotted is datafrom TiAlN (Derflinger et al., 1999) and TiN (Martin et al., 1999) whichshows surprisingly similar behaviour, suggesting that internal stresses mightaccount for a substantial fraction of the hardening observed earlier in Fig.8.2. This is consistent with the observation of differences in the hardnessesof monolithic materials and longer wavelength multilayer materials, by almosta factor of two. Furthermore the magnitude of such stresses is dependent on

10 20 30 40 50 60 70 80Hardness (GPa)

12

10

8

6

4

2

0

Inte

rnal

str

ess

(GP

a)

Multilayers

TiAlN/YNTiAlYN/VNTiAlN/VNTiAlN/CrNTiAlN/ZrNCrN/NbN

Monoliths

TiAlNTiN IAADTiN FAD

8.5 Variation of the internal stress and measured hardness for avariety of multilayer films compared with monolithic films. Themultilayer films are shown with filled symbols and the monolithicfilms with unfilled ones. The TiN was made by either ion assisted arcdeposition (IAAD) or filtered arc deposition (FAD). Note thathardnesses in the monolithic materials of up to 50 GPa appear to beachievable simply by internal stresses.


the processing conditions, consistent with the variability in the observedbehaviour. It is also consistent with the observation that heating can cause areduction in the hardness in TiN/NbN multilayers. This is normally attributedto a decrease in the sharpness of the interface (Barnett et al., 2003) but alsooccurs in monolithic films (Jindal et al., 1999).

8.4.2 Deformation processes and microstructureof the film

In uniformly strained materials, deformation structures can be readily observedusing transmission electron microscopy. However, it is much more difficultto prepare a similar sample where the deformation is more localized, as isthe case of nanoindentation. Recently this situation has been revolutionizedby the development of focused ion beam techniques for semiconductorprocessing, so that it is possible to select the region to be thinned to within100 nm (Overwijk et al., 1993; Saka, 1998).

Figure 8.6 shows an example of a cross-section made through ananoindentation in a TiN/NbN multilayer grown on an MgO substrate (Molina-Aldareguia et al., 2002). As expected, shear occurs on (101) and (10 1 )planes inclined at approximately 45∞ to the surface of the film (marked withblack arrows). However, other processes are also taking place. In particular,the crystal close to the indentation is deformed by rotation of the latticeplanes and has been observed elsewhere both in multilayers and in monolithicmaterials (Hultman et al., 1999). Although the situation here is not yetunderstood, such rotations can occur by the development of an array ofgeometrically necessary dislocations, associated with the strain gradientsunder the indent (Fleck et al., 1994) and which have in themselves beenassociated with hardening (Ashby et al., 1989). In Fig. 8.6(a) it can be seenthat flow under an indentation can occur by dislocation motion across theplanes, whereas in Fig. 8.6(b) it occurs by compression of the layers: comparepoints A and B.

Another feature that can be seen in Fig. 8.6(a) is the shearing that hasoccurred in the columnar grain boundaries in the film (marked with thewhite arrows). Sometimes these boundaries can grow through the completethickness of the film and lead to a section of the film being punched out, witha consequent reduction in hardness. This weakness is associated with thepores that are found along these boundaries. Such effects are well known inmonolithic films, where their number and size can be diminished by increasingthe surface mobility of adatoms on the growing film, for instance by raisingthe substrate temperature or increasing the flux of bombarding ions (Esteand Westwood, 1987; Thornton and Hoffman, 1989; Hoffman, 1994). However,the pore distribution can also be influenced by the multilayer structure, asshown in Fig. 8.7, which shows how the size and number of the pores is


Latticerotations

200 nmMgO (001)

(a)

(b)

25 nm

Bottom ofindent

AC

B

C

8.6 Cross-sectional TEM images of a 10 mN indent in a TiN/NbNmultilayer. White arrows point to a pre-existing columnar boundaryalong which shear can be observed. Black arrows point toconcentrated shear along {101} crystallographic planes; (b) shows ahigh-magnification image under the tip of the indenter, where it canbe seen that deformation has occurred by the layers at point Bhaving decreased compared to those at A.


reduced as the bilayer thickness, L, decreases, giving an increase in thehardness. Molina-Aldareguia has also shown how the structure of TiN in amultilayer adopts the denser structure of the NbN, again suggesting thatimprovements in the layer structure might be important (Molina-Aldareguia,2002).

A rather extreme example of this is given by Wang et al. in a study of TiN/AlN multilayers (Wang et al., 1998). For samples made with pulsed d.c.substrate bias, the improved hardness was retained down to the lowest valuesof L. However, when r.f. substrate bias was used the structure at L < 2 nm

8.7 Cross-sectional TEM images of (a) a CrN–AlN multilayer (the AlNis the lighter phase) and (b) monolithic CrN. Note that there are longpores present in the monolithic CrN, marked by the black arrow, butthese are not present in the multilayer.

(a)

(b)


was made up of columnar grains that had grown separately from theirneighbours and led to a decrease in hardness from 20 to 12 GPa. Changes inhardness at low values of L in this system have also been obtained elsewhere(Setoyama et al., 1996). The coating appears to behave rather like a series abed-springs where only that material directly under the indenter is pusheddownwards into the substrate, with no lateral constraint as would occur if thefilm were intact. Whilst this can be envisaged for the case where the grainsare completely separate, it cannot occur where the boundaries are porous butconnected, as is more commonly observed (Hultman et al., 1992; Molina-Aldareguia et al., 2000, 2002). In this case porosity presumably reduces thestress at which a columnar boundary will fracture, giving rise to a segmentof multilayer being pushed into the film. However, there is no quantitativedescription of such a process or of the microstructural variables that influenceit, despite its importance in limiting the hardness of the film.

This increase in hardness has been associated with the formation of arocksalt cubic (c) form of AlN, stabilized by the reduction of the interfacialenergy with the rocksalt-structured TiN (Madan et al., 1997). As the thicknessincreases the effect is offset by the increase in the volume free energy, so thatthe cubic form is stable only at layer thicknesses of less than 2 nm, althoughsome increase is possible by using rocksalt-structured compounds such asVN with a lower mismatch (Li et al., 2002, 2004). Stabilization with othermaterials such as W (Kim et al., 2001) or ZrN (Wong et al., 2000) is alsopossible, and other compounds such as CrN (Yashar et al., 1998) can alsoshow stabilized cubic forms.

The reason for the increase in hardness has been attributed to the bulkmodulus of c-AlN being greater than that of the wurtzite (w) structure. Thebulk modulus of c-AlN has been predicted to be 270 GPa compared with themeasured value of 205 GPa for w-AlN (Christensen and Gorczyca, 1993).The effect is two-fold. The higher bulk modulus is likely to give rise to ahigher hardness as well as an increased elastic mismatch, assuming thePoisson ratio is similar in both crystal structures. However, as TiN has a bulkmodulus greater than both forms of AlN, it is clear that any transformationto the wurtzite structure will lead to a reduction rather than an increase inrepulsive stresses due to any modulus mismatch.

The lattice resistance of the two crystal forms of AlN is not known.However, some tentative conclusions can be made by assuming that thelattice resistance at room temperature is close to the Peierls stress (Peierls,1940), which agrees with experimental observations of a wide range ofmaterials to within a factor of 3 and is given by

tn

pn

p = 2

1 – exp

21 –

G

db

◊ÊË

ˆ¯ (8.13)

where d is the spacing of the atoms across the slip plane and b is Burgers’


vector. B1 rocksalt-structured TiN shears on the {110}<110> slip system atroom temperature. If cubic AlN slips on the same slip system at roomtemperature as TiN, then the ratio d/b will be 0.5. In w-AlN, like others withthe same structure, slip takes place by the movement of partial dislocationson the more closely spaced {0001} planes (Delavignette et al., 1961; Yonenaga,2002), the glide planes, where d/b is 0.354. This suggests that tp /G for thelatter should be almost 10 times that of the cubic form of AlN, so that despitethe increase in the elastic modulus due to the increased packing density, w-AlN would be expected to show a higher lattice resistance and hence beharder than the cubic form. At best the differences are minimal. This suggeststhat the increase in hardness where L < 2 nm was due to the structure beingepitaxial, whereas at higher L it was polycrystalline, being made up of grainsof both c- and w-AlN.

8.4.3 Summary

In summary, it is clear that there are substantial effects that vary systematicallywith the wavelength of the multilayer due both to internal stresses and themicrostructure of the coatings. It has also been seen that deformation canoccur not just by dislocation flow, as the initial analyses have assumed, butby mechanisms such as lattice rotations and shear along column boundaries.In addition, the use of indentation complicates the deformation field, so thatthe assumption that equal strains in both layers are required need not becorrect. These effects all influence the hardness but have not so far beenincluded in analyses.

8.5 Conclusions

It can be seen that ceramic multilayer structures have been produced withincrements of the hardness of up to 60 GPa, increasing the hardness by up toa factor of almost 3. Initial work in this area has developed a number ofideas, such as the effect of modulus mismatch, which in some cases givegood agreement with the models suggested but in many others do not. It issuggested that at least some of this discrepancy can be accounted for bydifferences in the microstructure and residual stress-state of the film, both ofwhich are often poorly characterized. Furthermore there is very little directevidence about how these structures deform and in particular about howdifferent layers must be strained in order to accommodate the indenter whenit is pressed into the sample. Further advances in this area will require thegreater use of numerical techniques to analyse the complex stress and strainbehaviour under the indentation, coupled with the use of recently developedtechniques that allow the localized deformation behaviour to be observed indetail.


8.6 Future trends

Multilayer structures offer the potential for great increases in hardness and,although this has not been discussed here, also in wear resistance. The greatestdisadvantage of such structures is that the maximum in the hardness observedappears over a very narrow range of wavelengths. This is undoubtedly achallenge where structures with complex surface structures have to be coated.The best way forward is undoubtedly to understand how such hardeningarises and how the peak that is observed may be smoothed out. This will nodoubt require a complex combination of hardening mechanisms, as used incomplex alloy systems. Such an understanding will require the detailedcharacterization of the deformation processes occurring in indentation andwear and the correlation of these structures, initially, with the deformationpatterns that might be expected from continuum models of flow under anindenter, before developing more complex atomistic models.

8.7 Further reading

A list of references is given below which will allow the reader to investigateany points of detail. A good place to start is the collection of papers in theMarch 2003 MRS Bulletin (Volume 28, issue 3), which provides a snapshotof some of the more recent work on hard materials. For a general study ofthin films, try M. Ohring, The Materials Science of Thin Films, 1992, publishedby Harcourt Brace Jovanovich, Boston, MA.

8.8 References

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Barnett, S.A., Madan, A., Kim, I. and Martin, K. (2003), ‘Stability of nanometer-thicklayers in hard coatings’, MRS Bulletin, 28, 169–172.

Benlahsen, M., Lepinoux, L. and Grilhe, J. (1993), ‘Image forces on dislocations: theelastic modulus effect’, Materials Science and Engineering, A164, 428–432.

Bolshakov, A., Oliver, W.C. and Pharr, G.M. (1996), ‘Influences of stress on the measurementof mechanical properties using nanoindentation: Part II. Finite element simulations’,Journal of Materials Research, 11, 760–768.

Cahn, J.W. (1963), ‘Hardening by spinodal decomposition’, Acta Metallurgica, 11, 1275–1282.

Cheng, Y.T. and Cheng, C.M. (2000), ‘What is indentation hardness?’, Surface andCoatings Technology, 133–134, 417–424.

Christensen, N.E. and Gorczyca, I. (1993), ‘Calculated structural phase-transitions ofaluminium nitride under pressure’, Physics Review B, 47, 4307–4314.

Chu, X. and Barnett, S.A. (1995), ‘Model of superlattice yield stress and hardnessenhancements’, Journal of Applied Physics, 77, 4403–4411.

Delavignette, P., Kirkpatrick, H.B. and Amelinckx, S. (1961), ‘Dislocations and stackingfaults in aluminium nitride’, Journal of Applied Physics, 32, 1098–1100.


Derflinger, V., Brandle, H. and Zimmermann, H. (1999), ‘New hard/lubricant coating fordry machining’, Surface and Coatings Technology, 113, 286–292.

Este, G. and Westwood, W.D. (1987), ‘Stress control in reactively sputtered AlN and TiNfilms’, Journal of Vacuum Science and Technology A, 5, 1892–1897.

Fleck, N.A., Muller, G.M., Ashby, M.F. and Hutchinson, J.W. (1994), ‘Strain gradientplasticity: theory and experiment’, Acta Metallurgica et Materialia, 42, 475–487.

Gil-Sevillano, J. (1979), ‘On the yield and flow stress of lamellar pearlite’, in Strength ofMetals and Alloys, Vol. 2 (ed. Haasen, P., Gerold, V. and Kostorz, G.), PergamonPress, pp. 819–824.

Helmersson, U., Todorova, S., Barnett, S.A., Sundgren, J.-E., Market, L.C. and Greene,J.E. (1987), ‘Growth of single-crystal TiN/VN strained-layer superlattices with extremelyhigh mechanical hardness’, Journal of Applied Physics, 62, 481–484.

Hoffman, D.W. (1994), ‘Perspective on stresses in magnetron-sputtered thin films’, Journalof Vacuum Science and Technology A, 12, 953–961.

Högberg, H., Birch, J., Oden, M., Malm, J.O., Hultman, L. and Jansson, U. (2001),‘Growth, structure and mechanical properties of transition metal carbide superlattices’,Journal of Materials Research, 16, 1301–1310.

Hubbard, K.M., Jervis, T.M., Mirkarimi, P.B. and Barnett, S.A. (1992), ‘Mechanicalproperties of epitaxial TiN/(V0.6Nb0.4)N superlattices measured by indentation’, Journalof Applied Physics, 72, 4466–4468.

Hultman, L., Wallenberg, L.R., Shinn, M. and Barnett, S.A. (1992), ‘Formation of polyhedralvoids at surface cusps during growth of epitaxial TiN/NbN superlattice and alloyfilms’, Journal of Vacuum Science and Technology A, 10, 1618–1624.

Hultman, L., Engström, C., Birch, J., Johansson, M.P., Odén, M., Karlsson, L. andLjungcrantz, H. (1999), ‘Review of the thermal and mechanical stability of TiN-basedthin films’, Zeitschrift für Metallkunde, 90, 803–813.

Jayaweera, N.B., Downes, J.R., Frogley, M.D., Hopkinson, M., Bushby, A.J., Kidd, P.,Kelly, A. and Dunstan, D.J. (2003), ‘The onset of plasticity in nanoscale contactloading’, Proceedings of the Royal Society London A, 459, 2049–2068.

Jindal, P.C., Santhanam, A.T., Schleinkofer, U. and Schuster, A.F. (1999), ‘Performanceof PVD TiN, TiCN and TiAlN coated cemented carbide tools in turning’, InternationalJournal of Refractory Metals and Hard Materials, 17, 163–170.

Johnson, K.L. (1970), ‘The correlation of indentation experiments’, Journal of Mechanicsand Physics of Solids, 18, 115–126.

Kato, M., Mori, T. and Schwartz, L.H. (1980), ‘Hardening by spinodal modulated structure’,Acta Metallurgica, 28, 285–290.

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Kelly, A., Groves, G.W. and Kidd, P. (2000), Crystallography and Crystal Defects, JohnWiley & Sons, Chichester.

Kendall, K. (1975), ‘The effects of shrinkage on interfacial cracking in a bonded laminate’,Journal of Physics D: Applied Physics, 8, 1722–1732.

Kim, I.W., Madan, A., Gunz, M.W., Dravid, V.P. and Barnett, S.A. (2001), ‘Stabilizationof zinc-blende cubic AlN in AlN/W superlattices’, Journal of Vacuum Science andTechnology A, 19, 2069–2073.

Kim, J.O., Achenbach, J.D., Mirkarimi, P.O., Shinn, M. and Barnett, S.A, (1992), ‘Elasticconstants of single-crystal transition-metal nitride, films measured by line-focus scousticmicroscopy’, Journal of Applied Physics, Vol. 72 5, 1805–1811.

Koehler, J.S. (1970), ‘Attempt to design a strong solid’, Physics Review B, 2, 547–551.


Krzanowski, J.E. (1991), ‘The effect of composition profile shape on the strength ofmetallic multilayer structures’, Scripta Materialia, 25, 1465–1470.

Krzanowski, J.E. (1992), In Materials Research Society Symposium Proceedings, Vol.239 (ed. Nix, W.D., Bravman, J.C., Arzt, E. and Freund, L.B.), Materials ResearchSociety, Boston, MA, pp. 509–515.

Lehoczky, S.L. (1978a), ‘Retardation of dislocation generation and motion in thin-layeredmetal laminates’, Physical Review Letters, 41, 1814–1818.

Lehoczky, S.L. (1978b), ‘Strength enhancement in thin-layered Al–Cu laminates’, Journalof Applied Physics, 49, 5479–5485.

Lewis, D.B., Wadsworth, I., Münz, W.-D., Kuzel, R. and Valvoda, V. (1999), ‘Structureand stress of TiAlN/CrN superlattice coatings as a function of CrN layer thickness’,Surface and Coatings Technology, 116–119, 284–291.

Li, G., Lao, J., Tian, J., Han, Z. and Gu, M. (2004), ‘Coherent growth and mechanicalproperties of AlN/VN multilayers’, Journal of Applied Physics, 95, 92–96.

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Ljungcrantz, H., Engström, C., Hultman, L., Olsson, M., Chu, X., Wong, M.S. and Sproul,W.D. (1998), ‘Nanoindentation hardness, abrasive wear, and microstructure of TiN/NbN polycrystalline nanostructured multilayer films grown by reactive magnetronsputtering’, Journal of Vacuum Science and Technology A, 16, 3104–3113.

Madan, A., Kim, I.W., Cheng, S.C., Yashar, P., Dravid, V.P. and Barnett, S.A. (1997),‘Stabilization of cubic AlN in epitaxial AlN/TiN superlattices’, Physical Review Letters,78, 1743–1746.

Madan, A., Wang, Y., Barnett, S.A., Engström, C., Ljungcrantz, H., Hultman, L. andGrimsditch, M. (1998), ‘Enhanced mechanical hardness in epitaxial nonisostructuralMo/NbN and W/NbN superlattices’, Journal of Applied Physics, 84, 776–785.

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Mirkarimi, P.B., Hultman, L. and Barnett, S.A. (1990), ‘Enhanced hardness in lattice-matched single-crystal TiN/V0.6Nb0.4N superlattices’, Applied Physics Letters, 57,2654–2656.

Molina-Aldareguia, J.M. (2002), ‘Processing and nanoindentation behaviour of nitridemultilayers’, in Department of Materials Science and Metallurgy, University ofCambridge, Cambridge, UK.

Molina-Aldareguia, J.M., Lloyd, S.J., Barber, Z.H., Blamire, M.G. and Clegg, W.J. (2000),In Materials Research Society Symposium Proceedings: Thin Film Stresses andMechanical Properties VIII, Vol. 594 (ed. Vinci, R., Kraft, O., Moody, N., Besser, P.and Shaffer, E. II), Materials Research Society, Boston, MA, pp. 9–14.

Molina-Aldareguia, J.M., Lloyd, S.J., Odén, M., Joelsson, T., Hultman, L. and Clegg,W.J. (2002), ‘Deformation structures under indentations in TiN/NbN single-crystalmultilayers deposited by magnetron sputtering at different bombarding ion energies’,Philosophical Magazine A, 82, 1983–1992.

Münz, W.-D., Lewis, D.B., Hovsepian, P.E., Schönjahn, C., Ehiasarian, A. and Smith, I.J.


(2001), ‘Industrial scale manufactured superlattice hard PVD coatings’, SurfaceEngineering, 17, 15–27.

Overwijk, M.H.F., Van der Hauvel, F.C. and Bulle-Lieuwma, C.W.T. (1993), ‘Novelscheme for the preparation of transmission electron microscopy specimens with afocused ion beam’, Journal of Vacuum Science and Technology A, 11, 202.

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Part III

Nanostructured ceramic composites

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243

9.1 Introduction

The application of ceramics has infiltrated almost all fields in the last 20years, because of their advantages over metals due to their strong ionic orcovalent bonding. But it is just this bonding nature of ceramics that directlyresults in their inherent brittleness and difficulty in machining. In otherwords, ceramics show hardly any macroscopic plasticity at room temperatureor at low temperatures like metals. Hence, superplasticity at room temperatureis a research objective for structural ceramics. In recent years, many researcheshave been carried out to investigate nanophase ceramic composites.

Depending on the matrix grain size, nanophase ceramic composites canbe classified in two fundamental groups. One is composed of micrometer-sized matrices dispersed with a nanometer second phase, which has attracteda lot of interest in the last 15 years. In this group, the second phase plays acrucial role that affects the microstructure and the properties. Niihara [1] hasdivided it into three types – intragranular, intergranular and intra-/intergranular– trying to relate the distribution of the second nanophase in the matrix.Niihara and co-workers have reported dramatic improvements in toughness,strength at room temperature and high temperatures, creep strength andthermal shock resistance by incorporating nanocrystalline dispersion in amicrocrystalline matrix.

The other group of nanophase ceramic composites is nanocrystalline matrixcomposites, also called nanoceramics, in which the matrix grain size is below100 nm. The nano–nano type microstructure will be formed when the secondphase is also nano-scaled. Nanoceramics exhibit promising properties due tothe changes in deformation mechanisms when the grain size is reduced to theorder of 100 nm. The superplasticity of nanocrystalline CaF2 and nanocrystallineTiO2 at low temperatures, reported by Karch et al. [2], indicates that ceramicsare learning to ‘bend’ instead of fracture. Furthermore, nanoceramics alsoshow high toughness, in which a novel toughening mechanism called

9Nanophase ceramic composites

L Y O N G L I, Beijing University of Technology, China


Ferroelectric Domain Switching is recognized, different from that in micro–nano type ceramic composites.

9.2 Micro–nano type ceramic composites

In early nanocomposites, hard and strong dispersoids, such as SiC, Si3N4,TiC, etc., were mainly incorporated into the matrix to improve the mechanicalproperties. But in later years, enhancement of fracture strength was alsoachieved by addition of even soft and weak dispersoids like metals, graphiteand h-BN [3–5]. The density, microstructure and mechanical properties ofnano-sized particulate dispersion nanocomposites were strongly dependenton the volume fraction of particulate dispersion and sintering conditions.

9.2.1 Hard nanoparticle dispersed nanocomposites

Hard nanoparticles usually have a higher sintering temperature than that ofthe matrix, so that the sintering temperature is increased with increasinghard-particulate content. In Al2O3/SiC systems, only 5 vol% SiC incorporationcan evidently cumber the densification process. The nearly full densitiesattained by hot-pressing (HP) were achieved at 1600∞C for 5 vol% SiC, at1700∞C for 11 vol% and at 1800∞C for up to 33 vol%, while 1400–1500∞Cwas needed for Al2O3 monolithic ceramics [1]. At the same time, matrixgrain growth was dramatically inhibited owing to the pinning action of dispersedparticles. The sintering temperature for Al2O3 is 1500∞C, in which the grainsize grew up to 2.6 mm with uneven distribution. Al2O3 grain size in 5 vol%and 10 vol% SiC dispersed composites are 1.6 mm and 1.4 mm, respectively,with homogeneous distribution, even though sintering temperatures reach1700∞C or higher [3].

In general, particles disperse according to their grain size and the varietyof the matrix. For Al2O3/SiC, finer particles disperse within the matrix grainsand larger particles at the grain boundaries. The critical grain size is typically200 nm. For the MgO/SiC, Al2O3/Si3N4 and natural mullite/SiC composites,nanoparticles homogeneously disperse within as well as at the grain boundaries,which were confirmed to be the intra/inter-type nanocomposites.

It was found that dramatic improvements in toughness, strength, creepstrength and thermoresistance could be achieved by incorporating nano-SiCdispersion in a microcrystalline matrix. Subsequently, similar improvementswere found in other nanocomposites [1].

Niihara [1] considered the improved toughness was mainly attributed tothe residual stress that results from differential thermal expansion coefficientsof two phases. In Al2O3/SiC systems, the tensile hoop stress, thought to beover 1000 MPa around the nanoparticles within the matrix grains, was generatedfrom the large thermal expansion mismatch. Thus, the material may be

Nanophase ceramic composites 245

toughened primarily by the crack deflection and bridging due to the nano-sized SiC particles within the matrix grains. However, it was argued by otherresearchers that the incorporation of SiC into Al2O3 works on toughness ina very limited degree. Hoffman et al. [6] and Ferroni and Pezzotti [7] foundthat in Al2O3/5 vol% SiC systems, cracks were not distinctly deflected orbridged by nano-sized SiC, and no R-curve existed in Al2O3/5 vol% SiCnanocomposites. Zhao et al. [8] considered that the dramatically improvedtoughening mentioned by Niihara is actually attributed to surface compressivestresses. After relief annealing at high temperature, the toughness of Al2O3/SiC nanocomposites consequently had a sudden decrease.

The remarkable refinement of matrix grains by the dispersions is associatedwith the sub-grain boundary formation in the matrix grains. Sub-grainboundaries were found to be formed due to the pinning and pile-up ofdislocations by intragranular hard particles, which were generated in thematrix during cooling from the sintering temperature by the highly localizedthermal stress within and/or around the hard particles caused from the thermalexpansion mismatch between the matrix and the dispersions. This thermalexpansion mismatch, on the other hand, causes residual compressive stressesat the matrix grain boundaries. Both effects strengthen the composites. Thesub-grain boundaries were more extensive for Al2O3/5 vol% SiC after annealingat 1300∞C, and then the fracture strength was further improved from ~1050MPa to 1550 MPa.

The improvement in high-temperature hardness and brittle–ductile transitiontemperature (BDTT) must be due to the pinning of dislocations by the nano-sized dispersions. The Al2O3/SiC, Al2O3/Si3N4 and MgO/SiC nanocompositesgive a notable improvement in high-temperature strength up to and over1000∞C. In particular, the greatest improvement in high-temperature strengthwas observed for the MgO/SiC nanocomposites [1]. It is well known thatgrain boundary sliding and/or cavitation are responsible for the high-temperature strength degradation of oxide ceramics. Thus, the enhancementin strength at high temperatures is mainly due to the prohibition of the grainboundary sliding or cavitation by the dispersions within the matrix grains.

9.2.2 Metal nanoparticle dispersed nanocomposites

The mechanical properties of ceramics were improved by the addition ofnano-sized metal particles that were dispersed in the ceramic matrix. Ceramic/metal nanocomposites consisted of an oxide ceramic and either refractorymetal such as in the Al2O3/W, Al2O3/Mo and ZrO2/Mo systems or a metalwith a low melting point such as in Al2O3/Ni, Al2O3/Cr, Al2O3/Co, Al2O3/Feand Al2O3/Cu systems. These composites were fabricated by hot-pressingfine ceramic and metal powder mixtures or by reducing and hot-pressing thematrix and metal oxide powders [9–15].


For refractory metal dispersed systems, an outstanding improvementin toughness has been obtained in Al2O3/5 vol% Mo – as high as7.1 MPa.m1/2. At the same time, the fracture strength was just 306 MPa,much lower than that of Al2O3 monolithic ceramics. In Al2O3/Monanocomposites, the strength decreased with increasing sintering temperature,while toughness was considerably improved with increasing Mo volumefractions or increasing sintering temperature. When sintered at low temperature,nano-sized Mo particles are dispersed within the matrix grains, which canimprove the strength of grain boundaries and induce transgranular fracture.With an increase of sintering temperature or Mo volume fraction, nano-sizedMo particles agglomerated and grew to elongated grains. These elongatedMo grains make cracks deflect and bridge, which plays a key role in toughnessimprovement. Plastic deformation of large Mo particles at the crack tip alsogives an important contribution to toughening.

In low melting point metal dispersed systems, the Al2O3/Ni system hasbeen studied to obtain the desired microstructure and improvement ofmechanical properties by modification of the microstructure. Moreover,considering the magnetic properties of the composites, it was expected toimprove both mechanical and magnetic properties by incorporating merelynanometer-sized Ni, Co, and Fe into an Al2O3 matrix.

The average grain size of Al2O3/5 vol% Ni composites is finer (0.64 mm)than that of monolithic Al2O3 (1.2 mm), which is due to the growth restraintby the homogeneous dispersion of fine nickel particles. Fine nickel particles,less than 100 nm in size, disperse homogeneously at the matrix grain boundaries,forming the intergranular nanocomposite.

The ferromagnetic properties of nickel, such as high coercive force, wereobserved because of the fine magnetic dispersions, which indicates a functionalvalue of structural composites. The strength was enhanced up to 1090 MPaby dispersing only 5 vol% of nickel, almost doubling the strength of hot-pressed monolithic Al2O3. The high strength value of the Al2O3/5 vol% Nicomposites could be explained by the refinement of the matrix grains, becauseof the growth restriction caused by fine nickel dispersion. With increasing Nicontent over 10 vol%, the strength decreased and attained a value of 700MPa for Al2O3/20vol%Ni. This variation could be caused by the agglomerationof Ni particles dispersion at higher contents.

9.2.3 Soft and weak nanoparticle dispersednanocomposites

The super-fine dispersoids with laminar structure and a low modulus couldbe expected to play a crucial role in improving the machinability of ceramicsand their mechanical properties. Among the various nanocomposites, h-BNreinforced composites showed excellent corrosion resistance to molten metal


and high thermo shock resistance as well as good machinability [5, 16–19].Vickers hardness of both micro- and nanocomposites decreased directly withan increase of BN content. The fracture strength was marginally higher thanthat of the monolithic counterparts by adding h-BN up to 5 vol% and thendecreased gradually with an increase in BN for incorporation of low strengthBN, caused by the aggregation of h-BN particles. Oku and co-workers [20]developed a chemical process to prepare nano-sized BN coatings on ceramicpowders by reducing boric acid and urea in hydrogen gas. Kusunose et al.[16] reported that Si3N4/BN nanocomposites with homogeneously dispersednano-sized h-BN in an amount of not less than 20 vol% possess both goodmachinability and outstanding strength as high as about 1100 MPa. To date,examples of such materials include Si3N4/BN, Sialon/BN, SiC/BN, Al2O3/BN and 3Y-ZrO2/BN.

The fracture strength of nanocomposites was considerably improved, incomparison to the corresponding microcomposites. This implies that a nano-BN coating on particle surfaces effectively inhibited the grain growth andavoided abnormal grains in the sintering procedure so that the nanocompositespossessed a fine microstructure. Although the strength of samples decreasedwith increasing BN content, the strength of samples containing 20 vol% BNremained at 487 MPa for Al2O3/BN and at 838 MPa for 3Y-ZrO2/BN.

From another point of view, BN exerted a different influence on thefracture toughness of Si3N4/BN, Al2O3/BN and 3Y-ZrO2/BN. Si3N4/BN istoughened mainly by b-Si3N4 rod-like grains and 3Y-ZrO2/BN by martensiticphase transformation. There is no other toughening mechanism expected forBN in Al2O3/BN. So it is easy to understand that by incorporating nano-sized BN the toughness of nanocomposites was marginally increased for theAl2O3/BN system and nearly the same for Si3N4/BN. Toughness of 3Y-ZrO2/BN nanocomposites decreased with increasing BN content but was higherthan that of conventional microcomposites. The decrease in the fracturetoughness is due to the spontaneous tetragonal–monoclinic transformationduring cooling of the composites from the fabrication temperature and thereforea decrease in the amount of the available transformable tetragonal zirconiapresent in the micro- and nanocomposites. However, the tetragonal–monoclinictransformation is relatively insensitive to BN content in the 3Y-ZrO2/BNnanocomposites. Although addition of BN no doubt introduces flaws andinfluences strength of the materials to some degree, the nanocomposites stillkeep more tetragonal ZrO2.

All the above nanocomposites showed improvements of thermal resistanceby h-BN addition. In particular, Si3N4/BN has a DTC (the temperature differenceabove which the residual strength decreases suddenly) as high as ~1500∞C,50% higher than that of Si3N4 monolithic ceramics.

It is noted that, as in hard nanoparticle dispersed systems, h-BN incorporationalso increases the consolidation temperature, and besides, there is a large


thermal expansion mismatch between the Al2O3 matrix and BN (great contrastdifference also exists in the ZrO2/BN system), which causes the tensile stressand cracks at the interfaces between the matrices and the weak BN grain.The sintered body with BN content over 20 vol% contained destructivecracks and hence resulted in a low bulk density and poor strength [17, 18].Therefore, compounds with a low coefficient of thermal expansion such asSi3N4 and SiC should be added into the system to obtain fully dense bodies.

9.3 Nano–nano type ceramic composites

Nanocrystalline materials form an exciting area of materials research becausebulk materials with grain size less than 100 nm have properties not seen intheir microcrystalline counterparts. But the brittleness of nanoceramics haslimited their potential for use in structural applications, namely, research onnanoceramics shows that they are not inherently tougher than theirmicrocrystalline counterparts. Many strategies have been proposed to improvethe mechanical properties of nanoceramics by using reinforcement by asecond-phase addition and hybridization to develop nanocrystalline matrixcomposite materials.

9.3.1 Variety of hardness

Siegel et al. [21] pointed out that the shift in mechanical properties of metalsand ceramics occurs in opposite directions as grains become nanocrystalline.At room temperature, the nanocrystalline metals are harder than coarse-grained metals because the dislocation movement is pinned by the grainboundaries. On the other hand, nanoceramics are softer than their micrometercounterparts because of possible grain boundary movement. Hardness ofnanoceramics complies with an inverted Hall–Petch model. That is, hardnessdecreases with a reduced grain size. For example, the average room-temperaturehardness value was 4.45 GPa for Al2O3/ZrO2 nanoceramics by HIP [22],which is one quarter of the value of a comparable conventional material. Butinterestingly, SPSed Al2O3/ZrO2 nanoceramics have a considerably higherhardness, 15.2 GPa [23]. There is still some controversy regarding the hardnessof nanoceramics in some cases, i.e., hardness increases with reduced grainsize, indicating a positive Hall–Petch relation. Hardness of TiO2 nanoceramics(100∞C) reaches 12.75 GPa, much higher than the 1.96 GPa for traditionalTiO2 ceramics [24].

9.3.2 Superplastic deformation at low temperature

Since the early report of superplasticity in a ceramic material in 1986, avariety of such materials have been shown to exhibit superplastic behavior


at high temperatures, usually close to their original sintering temperature.Plastic deformation of the yttrium-stabilized tetragonal zirconia polycrystal(Y-TZP) containing micro- or nano-scaled grain was systematically studied,which is consistent with a grain boundary sliding (GBS) mechanism. Similarstudies on fully stabilized cubic ZrO2 (c-ZrO2) single crystals have receivedconsiderable attention in recent years, because the slip planes and directionsof c-ZrO2 are readily identified. Creep behavior of monoclinic ZrO2 (m-ZrO2) was also studied. These studies were all done at high temperatures.

It is well established that the plastic deformation of crystalline solidsoccurs by the movement of lattice dislocations and/or diffusional creep. Therate of diffusion is expressed as

e s d = /( )b3B D d KTW (9.1)

According to the equation, the diffusional creep rate of a polycrystal may beenhance by reducing the crystal size d, and by increasing the boundarydiffusivity Db. Nanoceramics are therefore expected to exhibit enhanceddiffusional creep for two reasons: first, the reduction of the crystal size fromabout 10 mm to ~10 nm enhances the creep rate by a factor of 109, andsecond, the enhanced boundary diffusivity may increase the creep rate by~103, so that the total enhancement is ~1012.

Significant plastic deformation will occur by a GBS mechanism if apolycrystalline ceramic is generated with a crystal size of a few nanometers,without having examined whether individual grains had been deformed.Karch and co-workers [2] obtained in 1987 the nanocrystalline CaF2 andTiO2 (8 nm) by applying high pressure in a high vacuum environment aftercollecting the powder in a mold. They observed that conventionally brittleceramics became ductile, permitting large (~100%) plastic deformations at alow temperature (80∞C for CaF2 and 180∞C for TiO2) to follow the shape ofa corrugated iron piston under pressure. Based on the fact that the CaF2

cubic structure has many slip systems available, the ductility seems to originatefrom the diffusional flow of atoms along the intercrystalline interfaces.

It is noted, however, that both of the above CaF2 and TiO2 nanoceramicshad some amount of porosity. This may account for an apparent soft behaviorrelated to the superplastic deformation at low temperature, which does notyet reveal the plastic deformation characteristics in nanoceramics. Localizedsuperplastic deformation under cyclic tensile fatigue tests was observed byYan et al. on 3Y-TZP nanoceramics at room temperature [25]. Themicromechanism behind this phenomenon is argued to be essentially governedby grain-boundary diffusion. The contribution of dislocation slip might be inoperation as a parallel mechanism to develop slip band-like microfeatures.

It was shown that, after cyclic tensile fatigue fracture at room temperaturein a narrow region within a couple of micrometers of each side of the fracturedsurfaces, the morphology of grain elongation appeared to be general. Viewed


on the fractured surface, the ratio of the long and short axes of the grains wasin the range 4–5 to 8–10. When the imaging was manipulated to track downa few micrometers on the side surface, it could be seen that the originalequiaxed grain morphology remained. It was also observed by AFM imagingthat microfeatures such as slip bands were developed on the side face of the3Y-TZP nanoceramics. For comparison, the microstructure of the surfaces ofthe specimens with an average grain size of 0.35 mm, after fatigue failure,showed equiaxed grains that had retained their original morphology.

It is easy to believe that the grain-boundary diffusion mechanism was themajor one to be considered, as described in equation (9.1). The direct effectof fine grain size alone on the rate of deformation can be obtained from thee –3µ d relationship. Therefore the 100 nm sized 3Y-TZP material shouldexhibit some 40 times higher rates of deformation in comparison with the0.35 mm grain-sized ones, under similar conditions.

9.3.3 Superplasticity of Si3N4/SiC nanoceramics at hightemperature

Si3N4 and SiC are promising structural materials for mechanical applications,especially in forming wear-resistant components such as engine parts, becauseof their excellent mechanical properties such as strength and hardness. Unlikein ionic crystals, plastic deformation of covalent compounds such as Si3N4

and SiC by dislocation glide is difficult because of their high Peierls force.Superplasticity has also been observed in some ionic crystals, such as Y-ZrO2. Wang and Raj [26] pointed out that superplasticity might occur inliquid-phase-sintered Si3N4 by diffusional creep enhanced by solution andprecipitation of crystals in a Si–O–N liquid phase at the grain boundaries.High ductility in compression has been observed, because cavitation at grainboundaries under tension and subsequent fracture occurs readily in the presenceof an intergranular liquid phase. Wakai et al. [27] reported superplasticdeformation of a covalent crystal composite, Si3N4/SiC nano–nano ceramics,which could be elongated by more than 150% at 1600∞C.

The mechanism of grain-boundary sliding depends on the structure of thegrain boundary. A liquid phase present at grain boundaries will enhancesliding. The microstructural changes during deformation were as follows:(1) phase transformation from a-Si3N4 to b-Si3N4 and grain growth of b-Si3N4, (2) crystallization of the intergranular liquid phase Si3N4·Y2O3, and(3) volatilization of the liquid phase in the Y–Si–Al–O–N system, causingweight loss.

The constitutive equation for steady creep can generally be expressed as

e s = /A dn n (9.2)

From the relationship between flow stress at a true strain of 0.1 and the strain


rate, n ª 2 was obtained. Models of diffusional creep enhanced by a liquidphase predict n = 1, and thus the superplasticity observed cannot be explainedby such models. The superplasticity exhibited by these covalent polycrystalswas characterized by non-Newtonian flow. The liquid-phase-enhanced creepmodel must be modified to accommodate this fact. A stress exponent of n =2 has been observed in superplastic alloys that do not contain an amorphousphase, and in zirconia polycrystals (Y-TZP) with a very thin amorphous filmat two-grain junctions, as well as in their Si3N4/SiC composites. Thus non-Newtonian flow was considered to be a common feature of superplasticity infine-grained polycrystalline materials. The observed superplasticity wasconsidered to be related to the presence of an intergranular liquid phase.Combined with its hardness, this property suggests several useful applicationsfor the novel material, for example to form engine components. Superplasticitywill make it readily moldable at high temperatures. It is suggested that theapparent strain hardening can be attributed to these microstructural changes.

9.3.4 High toughness and its toughening mechanism

Al2O3 /CNT nanoceramics

Carbon nanotubes (CNTs) offer a kind of nano-sized reinforcement that islightweight, has a hollow core, and has immense aspect ratio. Both theoreticaland experimental studies showed that carbon nanotubes have exceptionallyhigh mechanical properties such as strength, stiffness and flexibility, as wellas electrical and thermal conductivity, i.e. CNTs have a Young’s modulusapproaching 1 GPa, and exceptional tensile strength, in the range 20–100GPa. Especially, single-wall carbon nanotubes (SWCN) have an expectedelongation to failure of 20–30%, which combined with the stiffness (Young’smodulus of ~1.5 TPa) predict a tensile strength well above 100 GPa. Theflexibility of SWCN is remarkable and the bending may be fully reversibleup to a critical angle as large as 110∞. Hence, CNTs are considered to be theutmost type of fiber-like reinforcements. The excellent toughness of CNTsshould be helpful in solving the inherent brittleness of ceramics. In recentyears, attempts have been made to develop advanced engineering materialswith improved or novel properties through the incorporation of CNTs inselected matrices of polymers, metals and ceramics.

The bridging effect on cracks and the pullout of CNTs from the matrix arepossible mechanisms leading to the improvement of the fracture toughness.The contribution to toughness from the nanotube bridging for crackingtransverse to the axis of the nanotubes was calculated to be similar to5 MPa.m1/2. From the CNT bridging law, the CNT strength and interfacialfrictional stress were estimated to range from 15 to 25 GPa and from 40 to200 MPa, respectively [28]. These preliminary results demonstrate that


nanotube-reinforced ceramics can exhibit the interfacial debonding/slidingand nanotube bridging necessary to induce nanoscale toughening, and suggestthe feasibility of engineering residual stresses, nanotube structure, andcomposite geometry to obtain high-toughness nanocomposites. But actually,the use of CNTs to reinforce ceramic composites has not been very successful,in comparison to conventional ceramic fibers. Real toughness was not observed,or was fairly limited; even in situ or surface modification of CNTs solvedtheir dispersion into the matrix. The fracture toughness for the alumina-based composite containing 10 vol% of the in situ multi-wall carbon nanotubes(MWNT) was 4.2 MPa.m1/2, giving an improvement of only 24% over thatof the monolithic alumina [29]. Sun et al. [30] described a simple colloidalprocessing method to modify the surface of CNTs. The addition of 0.1 wt%CNTs in the alumina composite increases the fracture toughness from 3.7 to4.9 MPa.m1/2, about a 32% improvement.

It is indicated that the toughness of CNT-contained nanocomposites isdirectly dependent on the variety of CNTs and the sintering process. On theone hand, there are differences in the ability to transfer load from the matrixto the nanotubes between SWCN and MWCN, in addition to the differenceof their mechanical properties. The internal shells of MWCN are unable tobond to the alumina matrix, and therefore tensile loads are carried entirelyby the external shell. On the other hand, to be effective as reinforcing elements,high-quality, undamaged CNTs must be effectively bonded to the matrix sothat they can actually carry the loads. Ceramic composites reinforced byCNTs consolidated by hot-pressing methods require higher temperatures andlonger duration. These sintering parameters must damage the CNTs in thecomposites, leading to a decrease or total loss of reinforcing effects. Forexample, in composites of CNTs plus metal and ceramic, some of the hot-pressing temperatures were as high as 1600∞C. With both the Al2O3 and theMgAl2O4 matrices, a fraction of the CNTs seems to be destroyed during thehot-pressing at 1500∞C; when using the MgO matrix, most CNTs are destroyedduring hot-pressing at 1600∞C [31]. It seems that the quantity of CNTsretained in the massive composite is more dependent on the treatmenttemperature than on the nature of the oxide matrix. CNT damage producesdisordered graphene layers which gather at matrix grain junctions, and matrixgrains grew up to submicrometer range without producing fully densenanocomposites.

Zhan et al. [32, 33] fabricated fully dense nanocomposites of SWCN witha nanocrystalline alumina matrix at sintering temperatures as low as 1150∞Cby spark-plasma sintering (SPS). The introduction of SWCN leads to refinementof grain size. Most of the a-Al2O3 grains were in the nanocrystalline range,around 200 nm. The fracture toughness of the Al2O3/5.7vol%SWCNnanocomposite is more than twice that of pure alumina and there is almostno decrease in hardness. A toughness of 9.7 MPa.ml/2, nearly three times that


of pure alumina, was achieved in the Al2O3/10vol%SWCN nanocompositewhen sintered under identical conditions.

Some interesting features of microstructure can be noted in Al2O3/SWCNnanocomposite. First, ropes of SWCN were distributed along grain boundariesto develop a network microstructure. Some nanotubes were entangled withalumina grains and some encapsulated nanoscale alumina grains. Second,intimate contact between SWCN and alumina was observed in this material,unlike in alumina nanocomposites reinforced by CNTs grown in situ wherethe cohesion between carbon nanotubes and the matrix was poor and pulloutsof carbon nanotubes were observed. Stronger bonding of ropes with thematrix can be seen in the fully dense nanocomposite, suggesting that theextent of interfacial bonding should be a factor in increasing the toughnessof the composites. Third, no other forms of carbon, such as graphite, weredetected along the grain boundary, indicating that the nanotubes were notdamaged during consolidation by SPS. The increase in the quality and quantityof SWCN may have resulted in easier transfer of the stress.

It should be noted that the new work showed that the toughness of thesenanocomposites can be severely overestimated when measured by the standardindentation method. For dense Al2O3/SWNT composites, Vickers (sharp)and Hertzian (blunt) indentation tests reveal that these composites are highlycontact-damage resistant, as shown by the lack of crack formation. However,direct toughness measurements, using the single-edge V-notch beam method,show that these composites are as brittle as dense Al2O3 (having a toughnessof 3.22 MPa.m1/2) [34]. This type of unusual mechanical behavior was alsoobserved in SPS-processed, dense Al2O3/graphite composites and other softdispersoids-contained systems such as Al2O3/BN composites.

Al2O3 / ZrO2 nanoceramics

Cottom and Mayo [35] claimed that no increase in toughness occurs in ZrO2

nanoceramics having density values close to theoretical unless the materialsare heat-treated such that grains become susceptible to phase transformation.Taking the cue from the literature on superplastic metals, it is preferable tohave a two phase in one at microstructure of nano size. The Al2O3/ZrO2

binary system was chosen as the candidate system. There are two reasons forthis choice. First, there is very little miscibility between the two phases, asper the phase diagram. Secondly, in coarse-grained materials, there is evidenceof enhanced stability of the grain boundary structure due to the presence oftwo phases. Recent studies have indicated that a different mechanism (ratherthan phase transformation) of toughening must be operative here. TheFerroelectric Domain Switching is responsible for the great increase oftoughness in the Al2O3/ZrO2 nanoceramics, which is now well established asa mechanism for enhanced toughness without undergoing transformation in


ZrO2. Physically speaking, ferroelasticity is similar to ferromagnetism orferroelectricity. Drawing the analogy, ferroelasticity can be characterized bythe existence of permanent strain and an energy dissipating hysteresis loopbetween the stress and strain axes. In such materials, new domains or twinscan be nucleated depending on the state of stress.

From the literature on superplastic deformation of metals, it was concludedthat nano–nano composites are preferred, in which the constituent phaseshave similar grain sizes. Indeed, fine-grained (submicron but larger thannanocrystalline) Al2O3/ZrO2 composites have been superplastically deformed.It was believed that these results in coarser than nanocrystalline grainedAl2O3/ZrO2 materials are applicable to nanoceramics as well.

The average value of HIP toughness in Al2O3/ZrO2 nanoceramics wascalculated to be 8.38 MPa.ml/2 [22]. A conventional ceramic material cracksup substantially under such a load. This means that the nano/nano compositeswere actually deforming plastically under load. Lange reported a toughnessvalue of 6.73 MPa.m1/2 for a conventionally processed material. Furthermore,the conventional material is referred to as zirconia toughened alumina, wherethe controlled martensitic transformation of metastable tetragonal zirconiato the stable monoclinic phase should lead to phase transformation andmicrocrack toughening. In the nano/nano case, the toughness measurementswere carried out after ensuring that zirconia was not of the type resultingfrom martensitic transformation by verifying with X-ray diffraction (XRD)after grinding the samples. It was determined that roughly 5% of zirconiawas transforming, but about 95% of the zirconia was stable, presumably dueto the nanocrystalline size as well as constraint by the alumina grains. Thus,the increase in toughness (compared to 4 MPa.ml/2 for pure alumina) has tobe ascribed to the incorporation of nanocrystalline zirconia in alumina, whichis also in a nanocrystalline size.

It is pointed out that Al2O3/ZrO2 seems to be a versatile system for obtainingtailored properties by designing the microstructure. For example, it is possibleto fabricate hybrid microstructures of nano–nano inter-type wheretransformation toughening may act in synergism with the present mechanism.Another strategy is to heat-treat the nano–nano composite such that the ZrO2

grains grow to such an extent that they transform during the propagation ofcracks. In this suggested processing approach, transformation tougheningwill be a major mechanism and the primary use of these composites will beat temperatures closer to room temperature. Yet another interesting applicationcan be at moderate temperatures (i.e., at somewhat lower than the processingtemperatures, e.g. 1100∞C) where there may be more ductility due to additionalgrain boundary sliding. The high temperature limit of application of suchmaterials will be somewhat more than 1200∞C, where there may be a ductileto brittle transition due to the growth of nanocrystalline grains to conventionalgrains.


A combination of very rapid sintering at a heating rate of 500∞C/min andat a sintering temperature as low as 1100∞C for 3 min by the spark plasmasintering (SPS) technique, and mechanical milling the starting g-A12O3

nanopowder via a high-energy ball-milling (HEBM) process, can also resultin a fully dense nanocrystalline alumina matrix ceramic nanocomposite. Thegrain sizes for the matrix and the toughening phase were 96 and 265 nm,respectively. In regard to toughening, a great improvement in fracture toughness(~8.9 MPa.m1/2) was observed in the fully dense nanocomposite. It wasnearly three times as tough as the pure nanocrystalline alumina (152 nm,3.03 MPa.m1/2) [23].

It should be pointed out that the XRD study does not indicate any phasetransformation occurring during the 24 h HEBM period even though it islonger than the reported minimal time for the complete transformation(10 h). It is very interesting to note that the width of XRD for high-energyball-milled g-Al2O3 nanopowder became much greater than that for the startingnanopowder without HEBM. The residual stress induced by HEBM is likelyto be responsible for this wider XRD peak. Moreover, HEBM can lead tohigh green density due to pore collapse from the high compressive and shearstresses during the milling.

9.4 Fabrication of nanoceramics

Research on processing fully dense bulk nanoceramics and nanocompositesis attracting more and more interest. There have been some low-cost buteffective processes to obtain nano-sized ceramic powder and nanocompositepowders, such as sol-gel, micro emulsion, auto ignition, co-deposition andhigh-energy ball-milling (HEBM). One of the principal problems is the inabilityto consolidate nanopowders to high relative density without grain growth.To obtain the dense bulk nanoceramics, it is essential to decrease eithersintering temperature or retaining time at the highest point, or both HP, HIP,high-pressure sintering and fast consolidation techniques such as microwavesintering and spark plasma sintering (SPS) have been employed.

Among these techniques, high-pressure sintering seems to be the bestway of obtaining fully dense nanoceramics at the present time. The applicationof high pressure over 1–8 GPa results in a decrease of the temperature‘window’ within which fully dense compacts can be obtained without graingrowth or with only very limited grain growth. For Al2O3 based nanoceramics,success in achieving such fine grain size can be mainly attributed to twofactors. Firstly, lower sintering temperature no doubt leads to lower coarsening.Secondly, g Æ a transformation before sintering is quite important. Thistransformation is known to be nucleation controlled, and formation of avermicular structure during transformation hampers sintering transformation.Also, the presence of the second phase helps to restrain the grain growth.


The average grain size of fully dense alumina specimens was less than 100nm. The average grain size of Al2O3/Nb was in the range 40–50 nm. This isthe finest fully dense alumina-matrix composite reported so far. However,high-pressure sintering is limited to small and simple samples due to thehigh-pressure requirement [36].

Fast consolidation techniques, such as microwave sintering and plasmaactivated sintering (PAS), can enhance sintering and reduce the time availablefor grain growth. The advanced consolidation technique used in the presentstudy to overcome this hurdle is SPS. SPS is a comparatively new technique.It allows very fast heating and cooling rates, very short holding times, andthe possibility of obtaining fully dense samples at comparatively low sinteringtemperatures, typically a few hundred degrees lower than in normal hot-pressing. Unlike the first-generation spark sintering and the second-generationPAS, SPS can result in better control of the microstructure and properties ofmaterials in terms of sintering temperature and time. It is a pressure-sinteringmethod based on high-temperature plasma (spark plasma) momentarilygenerated in the gaps between powder materials by electrical discharge atthe beginning of on–off DC pulsing. This energizing method can generatespark plasma, spark impact pressure, joule heating, and an electrical fielddiffusion effect. In this process, powders are loaded into a graphite die andare heated by passing an electric current through the assembly. These processeshave now been developed beyond the production of small objects with simpleshapes, as continuous production of compacts of complex geometry and ofpieces with diameter larger than 150 mm has been achieved. Despite the factthat a uniaxial pressure is applied, green bodies of complex geometry can beexposed to a ‘pseudo-isostatic’ pressure when embedded in free-flowingelectrically conducting particulates that act as a pressure-transmitting mediuminside the die. By designing the mold, a temperature gradient along the directcurrent can be obtained, which is advantageous to simultaneously consolidatingcomponents with different sintering temperatures [37].

By optimizing the sintering parameters, various oxide powders weresuccessfully consolidated with nano- or submicron grain sizes, e.g. ZnO,Al2O3, ZrO2, YAG, Si3N4, SiC, Sialon and BaTiO3 [24, 38, 39]. The relatedgrain growth factor is not more than 2. A fully dense Al2O3/3Y-TZP nanoceramicwas obtained using SPS when the heating rate was increased to 500∞C/minup to 1100∞C for 3 min. The mean grain size for the alumina matrix was assmall as 96 nm with nearly 100% theoretical density. A ZnO powder withparticle size of 25 nm can be consolidated to 98.5% density with a grain sizeof ~100 nm. The heating rate is from 200∞C up to 550∞C, holding for oneminute, while it is consolidated at a temperature as high as 900∞C by microwavesintering.


9.5 Conclusions and future trends

Materials scientists now have the confidence and skills to conceive anddevelop ceramic materials with microstructures custom designed as micro–nano systems to obtain and improve the combination of strength, toughness,hardness, high temperature resistance, corrosion resistance and temperaturecreep resistance according to their application requirement. In situ reactionor coating is the main way to incorporate and homogeneously dispersenanoparticles into the matrix. Further research into enhancing the toughness,serviceability and machinability of nanocomposites is still required.Additionally, multifunctional ceramics incorporating the addition of a nano-sized second phase, which integrates both strong mechanical properties andsome electrical, magnetic, optical and calorific functions, will attract moreinterest. New features of ceramics, such as machinability and superplasticity,have been observed for the nano–nano composites. Today’s challenge ishow to obtain dense nanoceramics in an effective way at a low cost.

9.6 References

1. Niihara, K., New design concept of structural ceramics: ceramic nanocomposites, J.Ceram. Soc. Japan, 1991, 99(10): 974.

2. Karch, J., Birringer, R., Gleiter, H., Ceramics ductile at low temperature, Nature,1987, 330: 556.

3. Gao, L., Jin, X.H., Zheng, S., Ceramic Nanocomposites, Beijing: Chemical EngineeringPubl., 2004.

4. Suganuma, K., Sasaki G., Fujita, T. et al., Mechanical properties and microstructuresof machinable silicon carbide, J. Mater. Sci., 1993, 28(5): 1175.

5. Mizutani, T., Kusunose, T., Sando, M. et al, Fabrication and properties of nano-sizedBN-particulate dispersed Sialon ceramics, Ceram. Eng. Sci. Proc., 1997, 18(4B):669.

6. Hoffman, M., Rodel, J., Sternitzke, M. et al., Fracture toughness and subcriticalcrack growth in alumina/silicon carbide ‘nanocomposites’, Fracture Mechanics ofCeramics, 1996, 12: 179.

7. Ferroni, L.P., Pezzotti, G., Evidence for bulk residual stress strengthening in Al2O3/SiC nanocomposites, J. Am. Ceram. Soc., 2002, 85(8): 2033.

8. Zhao, J.H., Stearns, L.C., Harmer M.P. et al., Mechanical behavior of alumina–silicon carbide ‘nanocomposites’, J. Am. Ceram. Soc., 1993, 76: 503.

9. Oh, S.T., Lee, J.S., Sekino, T. et al., Fabrication of Cu dispersed Al2O3 nanocompositesusing Al2O3/CuO and Al2O3/Cu-nitrate mixtures, Scripta Mater., 2001, 44(8–9):2117.

10. Nawa, M., Sekino, T., Niihara, K., Fabrication and mechanical behaviour of Al2O3/Mo nanocomposites, J. Mater. Sci., 1994, 29(12): 3185.

11. Nawa, M., Yamazaki, K., Sekino, T. et al., A new type of nanocomposite in tetragonalzirconia polycrystal–molybdenum system, Mater. Lett., 1994, 20(5–6): 299.

12. Sekino, T., Niihara, K., Microstructural characteristics and mechanical propertiesfor Al2O3/metal nanocomposites, Nanostruc Mater., 1995, 6(5–8): 663.


13. Sekino, T., Niihara, K., Fabrication and mechanical properties of fine-tungsten-dispersed alumina-based composites, J. Mater. Sci., 1997, 32(15): 3943.

14. Ji, Y., Yeomans, J.A., Processing and mechanical properties of Al2O3–5 vol% Crnanocomposites, J. Eur. Ceram. Soc., 2002, 22(12): 1927.

15. Sekino, T., Nakajima T., Satoru U. et al., Reduction and sintering of a nickel-dispersed-alumina composite and its properties, J. Am. Ceram. Soc., 1997, 80(5):1139.

16. Kusunose, T., Sekino, T., Choa, Y.H. et al., Fabrication and microstructure of siliconnitride/boron nitride nanocomposites, J. Am. Ceram. Soc., 2002, 85(11): 2678.

17. Li, Y.L., Qiao, G.J., Jin, Z.H., Machinable, Al2O3/BN composite ceramics withstrong mechanical properties, Mater. Res. Bull., 2002, 38(7): 1401.

18. Li, Y.L., Zhang, J.X., Qiao, G.J. et al., Fabrication and properties of machinable 3Y-ZrO2/BN nanocomposites, Mater. Sci. Eng. A, 2005, 397: 35.

19. Wang, X.D., Qiao, G.J., Jin, Z.H., Fabrication of machinable silicon carbide–boronnitride ceramic nanocomposites, J. Am. Ceram. Soc., 2004, 87(4): 565.

20. Oku, T., Hirano, T., Kuno, M. et al., Synthesis, atomic structures and properties ofcarbon and boron nitride fullerene materials, Mater. Sci. Eng. B, 2000, 74(1–3): 206.

21. Siegel, R.W., Chang, S.K., Ash, B.J. et al., Mechanical behavior of polymer andceramic matrix nanocomposites, Scripta Mater., 2001, 44: 2061.

22. Bhaduri, S., Bhaduri, S.B., Enhanced low temperature toughness of Al2O3–ZrO2

nano/nano composites, Nanostruc. Mater., 1997, 8(6): 775.23. Zhan, G.D., Kuntz, J., Wan, J. et al., A novel processing route to develop a dense

nanocrystalline alumina matrix (less than or equal to 100 nm) nanocomposite material,J. Am. Ceram. Soc., 2003, 86(1): 200.

24. Gao, L., Li, W., Nanoceramics, Beijing: Chemical Engineering Publ., 2002.25. Yan, D.S., Zheng, Y.S., Gao L. et al., Localized superplastic deformation of

nanocrystalline 3Y-TZP ceramics under cyclic tensile fatigue at ambient temperature,J. Mater. Sci., 1998, 33(10): 2719.

26. Wang, J.G., Raj, R., Mechanism of superplastic flow in a fine-grained ceramiccontaining some liquid phase, J. Am. Ceram. Soc., 1984, 67(6): 399.

27. Wakai, F., Kodama, Y., Sakaguchi, S. et al., A superplastic covalent crystal composite,Nature, 1990, 344: 421.

28. Xia, Z., Curtin, W.A., Sheldon, B.W., Fracture toughness of highly ordered carbonnanotube/alumina nanocomposites, J. Eng. Mater. Tech., 2004, 126(3): 238.

29. Chang, S., Doremus, R.H., Ajayan, P.M. et al., Processing and mechanical propertiesof C-nanotube reinforced alumina composites, Ceram. Eng. Sci. Proc., 2000, 21(3):653.

30. Sun, J., Gao, L., Li, W., Colloidal processing of carbon nanotube/alumina composites,Chem. Mater., 2002, 14(12): 5169.

31. Flahaut, E., Peigney, A., Laurent, C. et al., Carbon nanotube–metal-oxidenanocomposites: microstructure, electrical conductivity and mechanical properties,Acta Mater., 2000, 48: 3803.

32. Zhan, G.D., Kuntz, J.D., Wan, J. et al., Single-wall carbon nanotubes as attractivetoughening agents in alumina-based nanocomposites, Nature Mater, 2003, 2: 38.

33. Zhan, G.D., Kuntz, J.D., Wan, J. et al., Plasticity in nanomaterials, Mat. Res. Soc.Symp. Proc., 2003, 740: 41.

34. Wang, X.T., Padture, N.P., Tanaka, H. et al., Contact-damage-resistant ceramic/single-wall carbon nanotubes and ceramic/graphite composites, Nature Mater., 2004,3(8): 539.


35. Cottom, B.A., Mayo, M.J., Fracture toughness of nanocrystalline ZrO2–3 mol%Y2O3 determined by Vickers indentation, Scripta Mater., 1996, 34(5): 809.

36. Mishra, R.S., Mukherjee, A.K., Processing of high hardness–high toughness aluminamatrix nanocomposites, Mater. Sci. Eng. A, 2001, 301: 97.

37. Tokita, M., Mechanism of spark plasma sintering, Nyu Seramikkusu, 1997, 10: 43.38. Nygren, M., Shen, Z.J., On the preparation of bio-, nano- and structural ceramics

and composites by spark plasma sintering, Solid State Sciences, 2003, 5(1): 125.39. Shen, Z.J., Zhan, Z., Peng, H. et al., Formation of tough interlocking microstructures

in silicon nitride ceramics by dynamic ripening, Nature, 2002, 417: 266.

260

10.1 Introduction

Carbon is the most versatile element in the periodic table. Due to variousbond structures such as sp3, sp2, sp hybrids, and multiple pp–pp bonds, it canform one-, two-, and three-dimensionally bond-structured substances andprovide a wide range of applications.1 Carbon materials such as graphite,diamond, activated carbons, carbon fibers, and C–C composites have beenextensively investigated and used for many years. Since the discovery ofcarbon nanotubes in 1997, carbon materials have been newly focused asfrontier materials in various fields.2–15

However, carbon materials have a serious shortcoming. They are easilyoxidized above 530∞C in air. It is possible to protect graphite plates orcarbon fibers with SiC coating by CVD or pyrolysis of polymer containingSi and C.16–18 SiC is known as an effective material to prevent oxidation andcorrosion due to the strong covalent bond and the passive oxidation byforming a protective SiO2 layer on SiC.19–24

It is difficult, however, to coat fine carbon materials such as carbon nanotubesand fine diamond powders with SiC uniformly. The SiC coating on carbonnanotubes would improve not only the oxidation resistance, but poor adhesionwith the matrix when they are used as nano-reinforcements. Many researchersindicate that the improvement of the adhesion between carbon nanotubesand the matrix is a critical issue to improve the mechanical properties oftheir composites.25–27 The SiC coating is very useful for fine diamond particlesas well. Diamond is widely used for cutting, grinding, and polishing ofvarious materials; however, the graphitization of diamond by the reactionwith transition metals such as iron, cobalt, and nickel limits its applications.If diamond particles could be coated with an effective protective layer, theycould be used at high temperatures under oxidizing and corrosive environmentsand the tool life could be extended. New composite formation of diamondwith WC/Co would be possible.

10Nanostructured coatings on advanced

carbon materials

Y M O R I S A D A and Y M I Y A M O T OOsaka University, Japan

Nanostructured coatings on advanced carbon materials 261

In this chapter, a new and easy process for SiC coating on finecarbon materials is described28–30 and some applications of SiC-coateddiamond particles and carbon nanotubes to create new composites aredemonstrated.31–33

10.2 Coating method of nanostructured SiC

10.2.1 Coating assembly

The SiC coating is processed based on the reaction of SiO vapor and carbonmaterials. Commercial SiO powders (99.9% pure) are provided as the siliconsource. The carbon materials are placed on the SiO powder bed via a carbonfelt as illustrated in Fig. 10.1. This assembly is covered with carbon sheetsin an alumina crucible to keep the SiO gas pressure in the crucible, andheated in a vacuum furnace at various temperatures from 1150 to 1550∞C invacuum (about 0.03 Pa) for periods of time between 1 and 90 minutes. It isnecessary to heat at a temperature greater than 1150∞C for the vaporizationof solid SiO.

10.2.2 SiC Coating on diamond particles

Microstructure

Diamond powders with a particle size of 1–30 mm are used for the SiCcoating. Figure 10.2 shows a TEM image of a SiC-coated diamond with aparticle size of ~ 1 mm. Each diamond particle is completely covered with apolycrystalline SiC layer ~60 nm thick. The grain size of SiC is severalnanometers. Although a large thermal expansion mismatch exists betweenSiC (a = 4.6 ¥ 10–6/K) and diamond (a = 3.1 ¥ 10–6/K), no crack or debonding

Alumina crucibleGraphite cover

Carbon materials

SiO(s)Carbon sheet

Carbon felt

10.1 Assembly for the SiC coating of carbon materials.


occurs in the SiC layer or the interface. If there were a gradual change incomposition from diamond to SiC at the interface, the thermal stress wouldbe relaxed.

Growth mechanism of SiC

Figure 10.3 shows the relation between temperature and pressure in thefurnace when the assembly of SiO powders, carbon sheets and carbon feltsis heated. The increase of the total pressure in the furnace at about 1200∞Cresults from the vaporization of SiO according to reaction (10.1). The evolutionof CO gas by the formation of SiC according to reaction (10.2) causes theincrease of the total pressure after vaporization of SiO.34–36 In this case, the

200 nm

Diamond

b-SiC layer

10.2 TEM image of the SiC-coated diamond particle.

0 100 200Time (min)

0.04

0.03

0.02

0.01

Deg

ree

of

vacu

um

(P

a)

Tem

per

atu

re (∞C

)

1400

1000

600

200

10.3 Temperature and degree of vacuum in the furnace.


surface of carbon sheets and carbon felts should react and transform to SiC,producing CO gas. Then, the diamond surface reacts with SiO(g) and formsa thin SiC layer on diamond. This SiC layer will act as a protective layer tolimit reaction (10.2) to produce further, thus limiting the evolution of CO(g)from diamond.

SiO(s) Æ SiO(g) (10.1)

2C(s) + SiO(g) Æ SiC(s) + CO(g) (10.2)

The SiO vapor is consumed within 30 min under these treatment conditions.Further treatment over 30 min causes thinning of the SiC layer, probably dueto the active oxidation taking place according to reaction (10.3). The partialpressure of oxygen in the furnace is about 6.0 ¥ 10–3 Pa and this value at thecoating temperature belongs to the active-oxidation region.37–40 Oxygen iscontinuously supplied from the outer atmosphere.

SiC(s) + O2(g) Æ SiO(g) + CO(g) (10.3)

The SiC-coated diamond particles can be characterized by X-ray powderdiffractometry. The diffraction peak appears at 35.6∞ which is assigned asthe b-SiC (111) plane.

The mechanism of SiC coating on diamond can be analyzed as follows.When the SiC-coated diamond particles are placed in an alumina containerand heated to 1200∞C in an airflow using a thermogravimetric apparatus, thesample weight decreases with increasing temperature mainly due to theoxidation of diamond, and reaches a minimum at about 1000∞C where diamondis almost completely converted to CO2 gas, leaving SiC behind. When it isfurther heated above 1000∞C, the minimum weight increases slightly due tothe passive-oxidation of the SiC layer. Because the mass gain due to thesilica formation on SiC below 1000∞C is negligibly small, we can determinethe minimum weight as the initial weight of the SiC layer on diamond. Letus consider a model structure consisting of a diamond sphere that is coateduniformly with SiC. Based on this model, the initial thickness of the SiClayer is expressed using the following equation:

W n r lf i2

SiC/ = 4p r (10.4)

where Wf is the minimum weight corresponding to the initial weight of theSiC layer, n is the number of diamond particles, ri is the average radius of adiamond particle, l is the thickness of the SiC layer, and rSiC is the densityof SiC (3.2 g/cm3).

The number of diamond particles can be obtained as follows:

n = (Wi – Wf)/(4/3)p ri3

diar (10.5)

where Wi is the initial weight of a SiC-coated diamond, and rdia is the densityof diamond (3.5 g/cm3).


By substituting equation (10.4) into (10.5), the thickness (l) of the SiClayer is calculated using the following equation:

l = Wfri rdia /3(Wi – Wf)rSiC (10.6)

Table 10.1 shows the calculated results on the thickness of the SiC layerand the mass gain due to the SiC formation depending on the coatingtemperature and time. The SiC coating for 90 min was obtained by repeatingthe 30 min coating three times. The thickness of the SiC layer increases withan increase in the coating time and temperature. The weight of the SiC layerincreases linearly with time. This result suggests that the growth of the SiClayer is not controlled by the self-diffusion of Si or C atoms through SiC, butby precipitation or deposition of SiC from the vapor-phase reaction. Thefollowing vapor–solid reactions account for the linear growth of the SiClayer with coating time:

SiO(g) + 3CO(g) Æ SiC(s) + 2CO2(g) (10.7)

CO2(g) + C(s) Æ 2CO(g) (10.8)

Based on these analyses on the SiC coating, the growth mechanism of theSiC layer on diamond is considered as follows. In the early stage of the SiCformation on diamond, a very thin SiC layer is formed on the diamondsurface according to reaction (10.2) between diamond and SiO(g). Once theSiC layer is formed, this reaction does not proceed due to the protectivelayer of SiC. The carbon sheet and felt in an alumina crucible act as thecarbon source. The reaction of CO2(g) with these carbon sources will producefurther CO(g) and deposit SiC(s) by reaction (10.7). Thin b-SiC whiskers areobserved on the surface of the SiC-coated diamond, suggesting the vaporgrowth of SiC.

The apparent activation energy of the SiC formation reaction is obtainedby an Arrhenius plot of the rate constants that can be calculated using themass gain data as a function of the coating temperature using the least-

Table 10.1 Thickness and mass gain of the SiC layer on diamond particle

Coating Coated at 1250∞C Coated at 1350∞C Coated at 1450∞Ctime (min)

Thickness Mass Thickness Mass Thickness Mass(nm) gain* (nm) gain* (nm) gain*

1 15 0.64 17 0.72 73 3.1315 21 0.89 29 1.24 82 3.5230 34 1.45 48 2.05 127 5.4690 79 3.47 109 4.68 235 10.18

*Mass gain is expressed by (DW/W0)106; W0 is the initial weight of diamond beforeSiC coating, and DW is the weight gain after coating.


squares method. The calculated value is 100 ± 21 kJ/mol. Shimoo et al.calculated the apparent activation energy for the formation of a SiC layer ona graphite plate based on reaction (10.7) and obtained 97 kJ/mol.36 Bothvalues show excellent agreement.

Figure 10.4 shows SEM photographs of the surface of SiC-coated diamondparticles coated at 1350∞C. Tiny granules of SiC were deposited and aggregatedwith an increase in coating time. Even for samples treated for 1 min, theentire surface is considered to be covered with a thin SiC layer formed by thedirect reaction of diamond and SiO(g) because the samples show good oxidationresistance, to be discussed later. EDX analysis shows a uniform distributionof Si atoms on the entire surface of the SiC-coated diamond particle.

Therefore, the SiC layer on diamond is considered to grow in a two-stepprocess as follows:

1. Formation of a very thin SiC layer by the direct reaction between SiO(g)and diamond.

2. Deposition of SiC on a thin SiC layer by reaction (10.7).

10.2.3 SiC Coating on carbon nanotubes

Microstructure

Multi-walled carbon nanotubes (MWCNTs) are coated with SiC becauseMWCNTs are more useful as reinforcements and more cost effective than

200 nm

200 nm

200 nm

200 nm

(a) (b)

(d)(c)

10.4 SEM images of the SiC-coated diamond treated at 1350 ∞C for(a) 1 min, (b) 15 min, (c) 30 min, and (d) 90 min.


single-walled carbon nanotubes (SWCNTs). The surfaces of MWCNTs arecovered with the SiC granules in the same way as the SiC-coated diamondparticles. The size of SiC granules is influenced by the coating temperatureand time. It is less than 50 nm for the sample treated at 1150∞C for 15 min,while it is about 150 nm for the sample treated at 1550∞C for 45 min.Therefore, the size of SiC granules can be controlled by adjusting the coatingconditions.

Figure 10.5 shows a high-resolution TEM photograph at the interfacebetween the SiC coating and MWCNTs, which was treated at 1350∞C for 15min. The (111) plane of b-SiC and the (002) plane of MWCNTs are clearlyobserved in the image. Some parts of MWCNTs at the vicinity of the interfacewith b-SiC show an amorphous structure. The measured angle between the(111) plane of b-SiC and the (002) plane of MWCNTs as shown in Fig. 10.5is 66–71∞. This angle matches closely the angle between two different (111)planes of b-SiC (70.5∞). These crystallographic relations suggest that the(111) plane of b-SiC is formed epitaxially on the (002) plane of MWCNTsand is grown toward the <111> direction.

Growth mechanism

Two types of assembly were used for the SiC coating to investigate thegrowth mechanism of the SiC layer. In the first method, the SiO powders areset on the bottom of an alumina crucible and MWCNTs are placed upon SiOpowders via a carbon felt, as shown in Fig. 10.6(a). In the second method, analumina plate with a center hole is used instead of the carbon felt to separatethe MWCNTs from SiO powders, as shown in Fig. 10.6(b). These assemblies

5 nm

SiC

MWCNTs3.54ÅCNTs (002)

2.57Åb-SiC (111)

10.5 High-resolution TEM image of the SiC-coated MWCNTs treatedat 1350 ∞C for 15 min with carbon source.


are covered with an alumina lid to keep the SiO gas pressure inside thecrucible, and heated at temperatures between 1250 and 1550∞C in a vacuumof about 0.03 Pa for 15 min and 30 min.

Figure 10.7(i) and (ii) show the XRD patterns of MWCNTs treated atvarious temperatures for 15 min. Figure 10.7(i) applies to the assembly ofFig. 10.6(a), and Fig. 10.7(ii) to Fig. 10.6(b). The diffraction peaks of b-SiCappear in all samples. For the samples treated in the assembly of Fig. 10.6(a),

Alumina lid

MWCNTs

SiO(s)

Alumina crucibleAlumina plate

Carbon felt

(a) (b)

10.6 Assemblies for the SiC coating of MWCNTs (a) with carbonsource, and (b) without carbon source.

: MWCNTs: b-SiC: a-SiC

(i) (iii)

(ii) (iv)

Inte

nsi

ty (

a. u

.)

Inte

nsi

ty (

a. u

.)

1550∞C

1450∞C

1350∞C

1250∞C

1550∞C

1450∞C

1350∞C

1250∞C

1550∞C

1450∞C

1350∞C

1250∞C

1450∞C

1350∞C

1250∞C

20 30 40 50 60 70 80Diffraction angle (2q/degree, CuKa)

20 30 40 50 60 70 80Diffraction angle (2q/degree, CuKa)

10.7 XRD patterns of the SiC-coated MWCNTs prepared (i) withcarbon source, and (ii) without carbon source for 15 min, and(iii) with carbon source, and (iv) without carbon source for 30 min,at various coating temperatures.


the peak of MWCNTs exists at 26.2∞(2q) which arises from the (002) graphitelayers. However, the peak intensity is very small for the sample treated at1450∞C in the assembly of Fig. 10.6(b). The peak of MWCNTs disappearswhen the sample is treated at 1550∞C. Above 1450∞C, the MWCNTs areconverted to SiC in the assembly of Fig. 10.6(b).

Figure 10.7(iii) and (iv) show XRD patterns for samples treated at varioustemperatures for 30 min. Figure 10.7(iii) applies to the assembly of Fig.10.6(a), and 10.7(iv) to Fig. 10.6(b). Both peaks of b-SiC and MWCNTsexist in the assembly of Fig. 10.6(a). However, the sample volatizes awaywhen treated at 1550∞C for 30 min in the assembly of Fig. 10.6(b). SiC isoxidized actively under a low oxygen potential at elevated temperature by areaction called ‘active oxidation’ following reaction (10.3). In this case, theoxidation occurs continuously and SiC is decomposed to SiO(g) and CO(g).The degree of vacuum in the furnace is ~0.03 Pa. This coating condition at1550∞C belongs to the active-oxidation region as reported by Schneider etal.37

Referring to Fig. 10.6, in assembly (a) the carbon felt would be oxidizedby the following reactions:

C(s) + O2(g) = CO2(g) (10.9)

C(s) + 12

O (g) = CO(g)2 (10.10)

A reducing atmosphere with very low oxygen content may not lead to asignificant active-oxidation reaction for SiC. Since the treatment for 30 minappears to be ineffective for forming a good SiC layer on MWCNTs, theSiC-coated samples prepared for 15 min were utilized to investigate thecoating mechanism.

There are some granules on the surfaces of MWCNTs coated in assembly(a). On the other hand, the surfaces of the SiC-coated MWCNTs prepared inassembly (b) are smooth. These morphological differences in SiC coatingssuggest that the formation process of the SiC layer is different. Reaction(10.7) would proceed when there is a rapid decrease in the partial pressureof CO2(g). Because there is carbon felt in the crucible in assembly (a),CO2(g) is converted to CO(g) by reaction (10.8). The CO(g) generated bythis reaction will be supplied for reaction (10.7). In assembly (b), it is difficultto promote reaction (10.8) because no extra carbon source exists in thecrucible. In this case, the surfaces of MWCNTs react directly with SiO(g)and convert to SiC by reaction (10.2).

From the above results, the growth model of the SiC formation on MWCNTscan be proposed as illustrated in Fig. 10.8. The growth of SiC is influencedby the existence of the carbon source in the crucible. In the early stage of thereaction, SiO(g) reaches the surface of MWCNTs and forms a thin SiC layer

Nanostructured coatings on advanced carbon m

aterials269

1 2 3

SiC formed by conversion (reaction �1 )

1 2 3

(a)

(b)

MWCNTs

MWCNTs

SiC formed by conversion (reaction �1 )

SiC formed by deposition(reaction �2 , �3 )

SiC formed by deposition(reaction �2 , �3 )

SiC formed by deposition(reaction �2 )

SiC formed by conversion (reaction �1 ) SiC formed by conversion (reaction �1 )

SiC formed by deposition(reaction �2 )

10.8 Growth models of SiC layer on MWCNTs: (a) with carbon source; (b) without carbon source.

(1) SiO(g) + MWCNTs Æ SiC(s) + CO(g)(2) SiO(g) + 3CO(g) Æ SiC(s) + 2CO2(g)(3) CO2(g) + C(s) Æ 2CO(g)


due to reaction (10.2). This stage needs no carbon source. In assembly (a),SiC granules are deposited on the thin SiC layer by reaction (10.7). Thisreaction continues until the SiO(s) is consumed because the partial pressureof CO2(g) is decreased by reaction (10.8). The generation of CO(g) by reaction(10.8) would control reaction (10.7). MWCNTs can remain after the SiCcoating in assembly (a).

In assembly (b), reaction (10.2) is promoted and MWCNTs are convertedto SiC. It is inferred that MWCNTs are changed to SiC nanorods by reaction(10.2) because the diffraction peaks of MWCNTs are not observed on XRDpatterns of the sample prepared at 1550∞C for 15 min. A small amount ofnanometer-scale SiC granules are deposited by reaction (10.7) even in assembly(b). In this case, the MWCNTs are considered to play the role of carbonsource.

10.2.4 Oxidation resistance

SiC-coated diamond particles

Table 10.2 shows starting temperatures of oxidation for SiC-coated diamondparticles depending on the coating time and temperature. The coating at1350∞C shows superior oxidation resistance. When the total coating time is90 min (30 min coating repeated three times), the starting temperature ofoxidation reaches about 950∞C, which is 400∞C higher than that of diamondwithout SiC coating. Even for diamond particles treated for only 1 min at1350∞C, no oxidation starts below 750∞C. This result suggests the existenceof a thin SiC layer on the entire surface of diamond. The thickness of thisSiC layer is estimated to be about 15 nm by interpolating the linear relationbetween the thickness of the SiC layer and the holding time at 1350∞C. Thediamond surface is converted to SiC by the reaction-diffusion of Si intodiamond. Other coatings at 1250∞C and 1450∞C show lower oxidationresistance. It is reported that the transformation of diamond to graphite andthe generation of cracks in diamond start at over 1400∞C.41 The coating at1250∞C for 90 min exhibits no improvement against oxidation compared tothe coating at the same temperature for 1 min.

Table 10.2 Starting temperature of oxidation for the SiC-coated diamondparticles

Coating Coated at Coated at Coated attime (min) 1250∞C (∞C) 1350∞C (∞C) 1450∞C (∞C)

1 747 742 79415 752 798 78630 753 848 80890 761 926 831


The oxidation durability of SiC-coated diamond particles treated at 1350∞Cfor 30 min and 90 min is evaluated at 700∞C in an airflow of 50 ml/min. Theuncoated diamond particles show a rapid weight loss, whereas the SiC-coated diamond particles treated for 90 min maintain over 70% of theirweight after oxidation for 5 h.

SiC-coated carbon nanotubes

It is easy to form SiC nanorods in the assembly shown in Fig. 10.6(b). TheSiC nanorod must show excellent oxidation resistance. However, themicrostructure and superior properties of MWCNTs are lost. Therefore, onlythe SiC-coated MWCNTs treated in the assembly shown in Fig. 10.6(a) for15 min are evaluated. Figure 10.9 shows the TG curves for SiC-coatedMWCNTs which are heated at 650∞C in air. As-received MWCNTs areoxidized completely within 5 min. The remaining mass detected (~2.5%) isattributed to iron because it is used as a catalyst for the synthesis of MWCNTs.Further heating increases the mass gain, probably due to the formation ofFe2O3. For the SiC-coated MWCNTs treated at 1550∞C, about 90% of massremains after heating at 650∞C for 60 min. The SiC coating at higher temperatureprovides an improved oxidation resistance.

The morphological change of MWCNTs before and after the oxidation at650∞C for 10 s is shown in Fig. 10.10(a) and (b). A tip of MWCNTs beforeoxidation is closed with a cap. On the other hand, the cap is removed after

as-received

1550∞C

1450∞C

1350∞C

1250∞C

0 10 20 30 40 50 60Holding time (min)

100

80

60

40

20

0

Mas

s lo

ss (

%)

10.9 TG curves for the SiC-coated MWCNTs heated at 650∞C in air.The coating was carried out at various temperatures from 1250∞C to1550∞C.


oxidation for 10 s. It is reported that the cap is not resistant to chemicalreactions because it has a pentagonal shape.42 When the cap is lost, oxygencan enter the interplanar spaces between [002] planes of MWCNTs. Suchplanes of carbon atoms are bonded by Van der Waals forces. It is well knownthat the (002) plane of graphite has a higher oxidation rate than other planes.Therefore, the tip of MWCNTs is very important to prevent oxidation ofMWCNTs. For the sample coated at 1350∞C for 15 min, the surface and capare covered with tiny SiC granules (Fig. 10.10(c)). The cap is held and about70% of the mass remains after oxidation at 650∞C for 1 h (Fig. 10.10(d)).

A relatively strong (002) peak of MWCNTs remains on the XRD patterneven after oxidation tests. The oxidation rate must be controlled by thediffusion of oxygen through the SiC layer. The crystalline size of SiC can beestimated using Scherrer’s equation for the (111) peak of b-SiC and is plottedas a function of coating temperature in Table 10.3. The crystalline size ofSiC gradually increases from 13 nm to 26 nm with an increase in coatingtemperature. The higher coating temperature can produce dense and thickSiC layers with larger crystalline size, resulting in a higher oxidation resistance.The shape of the SiC-coated MWCNTs is not changed by the oxidation test.No crack or exfoliation is observed on the surface.

1 mm

1 mm 1 mm

1 mm

(a)

(c)

(b)

(d)

10.10 SEM photographs of MWCNTs and SiC-coated MWCNTs:(a) as-received MWCNTs; (b): MWCNTs oxidized at 650∞C for 10 s;(c) SiC-coated MWCNTs; (d) SiC-coated MWCNTs oxidized at 650∞Cfor 60 min.


10.3 Applications of nanostructured SiC coatings in

advanced composites

10.3.1 Development of SiC-coated diamond/WC–Cocomposites

Cemented carbide, an alloy made of tungsten carbide (WC) and cobalt (Co),is widely used for cutting tools and wear-resistant tools because of its excellenthardness, strength, toughness, and Young’s modulus.43 Although sintereddiamond has extremely high hardness and wear resistance,43 it is costly andlimited in size and shape because of the need for using ultra-high pressureand difficulty in machining. Therefore, the composite formation of cementedcarbide and diamond under lower pressure is very attractive because newwear-resistant tools can be produced at lower cost. However, diamond reactswith cobalt at the sintering temperature of cemented carbide, ~1150∞C, andconverts to graphite. This reaction must be prevented to develop the diamond-dispersed cemented carbide.

Fabrication

SiC-coated diamond powders with a particle size of ~8–16 mm are mixedwith fine WC and cobalt powders (10 wt% Co) and sintered in a vacuum at1220∞C at 30 MPa for 5 min by pulsed-electric current sintering (PECS).44,45

The diamond content is 20 vol%. The PECS method enables sintering ofmaterials at a lower temperature and shorter time than the conventionalsintering methods because it uses a high pulsed current of 1000–3000 A.This current is sent through the material directly, after placing the materialin a graphite mold, under uniaxial loading.

Microstructure

Figure 10.11(a) is a SEM micrograph of the polished surface of a SiC-coateddiamond dispersed WC-10wt%Co composite. It is well sintered and the

Table 10.3 Relation between coating temperature andapparent crystalline size of SiC

Coating temperature Crystalline size(∞C) (nm)

1250 13.01350 19.41450 21.71550 26.0

Note: Coating time is 15 min.


diamond particles are uniformly dispersed. The relative density reaches 99.5%.The dense composite suggests that the SiC layer protects the diamond fromattack by molten cobalt and prevents the conversion of the diamond surfaceto graphite. In contrast, when uncoated diamond powders are mixed withWC–Co, well-sintered materials could not be obtained.


Values of Vickers hardness and indentation fracture toughness measured forthe WC-Co with and without SiC-coated diamond are compared in Table10.4. The Vickers hardness is measured under a 98 N load. The fracturetoughness is evaluated under the same load using the indentation fracturemethod. Both sintered composites show almost the same hardness of ~15.5GPa. However, the fracture toughness of the diamond-dispersed compositewas 16.3 MPa.m1/2, which is nearly 200% higher than that of WC–10wt%Coitself. Because of the extremely high hardness (~110 GPa) and high Young’smodulus (~950 GPa) of diamond, the presence of diamond particles is expected

(a)

(b)

100 mm

20 mm

10.11 SEM images of the SiC-coated diamond-dispersed cementedcarbide composite: (a) polished surface, (b) crack propagation.


to impede the crack propagation. Such an impeding effect against crackpropagation is seen in Fig. 10.11(b). The lower thermal expansion coefficient(~3 ¥ 10–6/K)46 than that of the cemented carbide (5.7 ¥ 10–6/K)47 and thehigh Young’s modulus of diamond would produce a high tensile stress aroundeach diamond particle. Such a high tensile stress can further enhance thecrack deflection.48 No increase in hardness of the cemented carbide compositewith SiC-coated diamond dispersion could be attributed to weak bondingbetween the diamond and cemented carbide matrix. The diamond-dispersedcemented carbides have a wear resistance ten times higher than that of theconventional cemented carbides. Such products are commercialized as superwear-resistant tools in Japan.49

10.3.2 Development of SiC-coated carbonnanotubes/WC–Co composites

MWCNTs have been tested to reinforce various matrices because they havemany unique mechanical and physical properties.14,15 However, these nanotubesbecome corroded with metals (such as iron, cobalt, and aluminum) attemperatures above 850∞C. These shortcomings limit the applications ofMWCNTs as nano-reinforcements. The SiC coating can effectively protectthe diamond from molten cobalt, thus allowing dense SiC-coated diamond-dispersed cemented carbide composites to be successfully fabricated at lowerpressures. If MWCNTs can be coated with the same SiC layer, more stableMWCNTs would be produced and expected to be used as nano-reinforcementsfor various matrices. The development of SiC-coated MWCNTs/WC-Cocomposites has potential to extend functions of both MWCNTs and WC–Co.

Fabrication

SiC-coated or uncoated MWCNTs are dispersed in isopropylalcohol (IPA)using ultrasonic vibration for 5 min. Then the SiC-coated or uncoated MWCNTsare mixed with fine WC and cobalt powders (10 wt% cobalt) in IPA usingplastic balls (10–20 mm in diameter) for 5 h. After ball milling, the IPAsolution is evaporated by stirring using an electric heater and then dried at

Table 10.4 Comparison of Vickers hardness and fracture toughness forcemented carbide sintered with and without SiC-coated diamond particles

Materials Vickers Indentation fracturehardness (GPa) toughness (MPa.m1/2)

WC + 10wt%Co 15.4 8.7WC + 10wt%Co + 20vol% 15.5 16.3SiC-coated diamond


100∞C. The MWCNTs are uniformly dispersed without aggregation. Themixed powders are sintered at different temperatures from 950∞C to 1200∞Cat 30 MPa for 5 min under vacuum using PECS. The content of the SiC-coated MWCNTs is 3 vol%.

Sintering behavior

The mixed powders of SiC-coated MWCNTs and WC–10wt%Co are chargedin a graphite mold and sintered by PECS. The sample is heated from roomtemperature to sintered temperatures for 20 min. Table 10.5 shows the changeof relative density of the WC–10wt%Co and the WC–10wt%Co with MWCNTscompacts depending on the sintering temperature. The density of WC–10wt%Co increases with an increase in sintering temperature and reachesnearly 100% at 1150∞C. On the other hand, the composites of WC–10wt%Cowith MWCNTs are fully sintered at 1050∞C, which is 100∞C lower than thesintering temperature of WC–10wt%Co itself. The resistance heating ofMWCNTs is capable of accelerating the sintering process.

Microhardness

The microhardness of the WC–10wt%Co compact increases by incorporatingSiC-coated MWCNTs except for the coating at above 1450∞C, as shown inTable 10.6. The microhardness is measured under a 19.6 N load. It is reasonable

Table 10.5 Relation between relative density of WC-10wt%Co and sinteringtemperature

Materials 1000∞C (%) 1050∞C(%) 1100∞C(%) 1150∞C(%)

WC + 10wt%Co 69.8 86.3 95.0 100.0WC + 10wt%Co + 3vol% 92.6 100.0 100.0 100.0SiC-coated MWCNTs

Table 10.6 Relation between microhardness of the WC-10wt%Co compactswith MWCNTs, SiC-coated MWCNTs, and without MWCNTs, and coatingtemperature

Materials Coating temperature Microhardness(∞C) (GPa)

WC + 10wt%CO — 17.5WC + 10wt%Co + MWCNTs — 18.9WC + 10wt%CO 1150 20.0+ SiC-coated MWCNTs 1250 20.0

1350 19.91450 18.31550 18.0


to assume that the hardness is enhanced probably by the elastic recovery ofMWCNTs. The increase of microhardness by incorporating uncoated MWCNTsis lower than that by SiC-coated MWCNTs. This difference of hardness maybe due to the corrosion of uncoated MWCNTs with molten Co and pooradhesion with the matrix as suggested by Laurent et al.25 The presence of theSiC coating overcomes these problems because SiC is chemically stable andthe SiC granules can provide an anchor effect. The lower hardness obtainedwhen SiC-coated MWCNTs prepared at over 1450∞C are used forreinforcements could be attributed to the strength degradation of MWCNTsas a result of high coating temperatures.

10.3.3 Development of SiC-coated carbonnanotubes/SiC composites

SiC has high heat and oxidation resistance. Therefore, various applicationsrelating to space developments and efficient power generators are expected.However, the low reliability due to the brittle nature of SiC is a criticalproblem. MWCNTs may be good candidates to reinforce the SiC matrix ifthe original strength of MWCNTs is maintained. The SiC coating is expectedto improve the weak adhesion between MWCNTs and the SiC matrix.

Fabrication

The SiC-coated or uncoated MWCNTs are dispersed in isopropyl alcohol(IPA) using ultrasonic vibration for 5 min. They are then mixed with nanometer-sized SiC (mean diameter 30 nm) and B4C (mean diameter 240 nm) powdersin IPA using ultrasonic vibration for 10 min. The B4C is added at 2 wt% asa sintering aid. After the mixing, the IPA solution is evaporated and themixed powders are dried at 100∞C. The MWCNTs are uniformly dispersedwithout aggregation. The mixed powders are sintered at 1800∞C at 40 MPafor 5 min under a vacuum by means of PECS. The content of the SiC-coatedMWCNTs is varied between 1 and 5 vol%.


The microhardness of the SiC compact measured under a 19.6 N load increasesby incorporating SiC-coated MWCNTs, as shown in Table 10.7. The hardnessreaches 30.6 GPa for the content of 5 vol% SiC-coated MWCNTs. This highhardness is considered as an apparent value due to the elastic recovery of theindentation after loading. This interesting phenomenon is discussed later. Onthe other hand, the increment of hardness by incorporating uncoated MWCNTsis very low compared with the increment by incorporating SiC-coatedMWCNTs. This behavior may be due to the poor adhesion with the matrix.


The SiC coating acts as an adhesive to the SiC matrix and the SiC granulesprovide an anchor effect. The relatively lower hardness is obtained whenSiC-coated MWCNTs prepared at over 1250∞C are used for reinforcements.It is believed that the high coating temperature causes the strength degradationof MWCNTs due to the conversion of MWCNTs to SiC. Table 10.8 showsthe values of fracture toughness measured under a 19.6 N load. The toughnessincreases to 5.4 MPa.m1/2 by the dispersion of SiC-coated MWCNTs, althoughthe data are quite scattered. The mean values of hardness and fracture toughnessof the compact measured under 9.8 N and 19.6 N loads are listed in Table10.9. The results are attributed to the improvement of the adhesion betweenthe MWCNTs/SiC matrix and the SiC coating. The hardness of the monolithicSiC and the uncoated MWCNTs dispersed SiC composite does not change inrelation to the indentation load, while that of the SiC-coated MWCNTs/SiCcomposite increases up to 34.3 GPa when the hardness test is carried out

Table 10.7 Relation between microhardness and MWCNT content

Content Without Coated at Coated at(vol%) coating (GPa) 1150∞C (GPa) 1250∞C (GPa)

1 26.7 25.7 26.43 25.4 29.0 27.15 25.8 30.5 27.4

Note: Microhardness of monolithic SiC is 25.5 GPa.

Table 10.8 Relation between fracture toughness and MWCNT content

Content Without Coated at Coated at(vol%) coating 1150∞C 1250∞C

(MPa.m1/2) (MPa.m1/2) (MPa.m1/2)

1 4.4 4.7 5.13 4.9 4.8 5.55 4.6 4.9 5.4

Note: Fracture toughness of monolithic SiC is 4.8 MPa.m1/2.

Table 10.9 Effect of indentation load on microhardness and fracture toughness

Materials Microhardness Fracture toughness(GPa) (MPa.m1/2)

<9.8 N> <19.6 N> <9.8 N> <19.6 N>

Monolithic SiC 26.1 25.5 4.0 4.8MWCNTs (5 vol%)/SiC 20.5 25.8 4.4 4.6SiC-coated MWCNTs 34.3 30.6 7.1 5.4(5 vol%)/SiC


under a 9.8 N load. The fracture toughness has a tendency to increase dependingon the decrease of the indentation load. These results suggest that uncoatedMWCNTs do not act as reinforcements for SiC matrix, due to weak interfacialadhesion between the surface of the MWCNTs/SiC matrix.

Figure 10.12 shows the SEM and three-dimensional (3D) images ofindentation prints marked by the hardness test under a 19.8 N load. The 3Dimages are composed on the same scale to compare the shape of indentationprints. These images are synthesized with signals of secondary electronsusing four detectors in the 3D-SEM equipment. The indentation print of themonolithic SiC ceramic is very sharp, reflecting a square pyramidal shape ofVickers hardness tester. The cracks propagate outward from each corner ofthe indent. The MWCNTs/SiC composite shows somewhat similar fractography.On the other hand, the indentation print and the crack propagation of theSiC-coated MWCNTs/SiC composite are very indistinct and the squarepyramidal print cannot be observed in the 3D image. However, the lateralsides of the indent indicate an elastic deformation. It is difficult to understandwhy such a high hardness of 34.3 GPa is obtained by incorporating only 5vol% of MWCNTs into the SiC matrix. If we suppose that the SiC-coatedMWCNTs are tough and exhibit an excellent load transformation effect, theelastic recovery of the indentation print would occur and show high hardness,apparently.

Microstructure

SEM images of the fractured surfaces of a monolithic SiC ceramic, uncoatedMWCNTs/SiC, and SiC-coated MWCNTs/SiC are shown in Fig. 10.13. Allsamples were dense and pore-free. For the uncoated MWCNTs/SiC composite,

(a) (b) (c)

10.12 SEM and 3D images of the indentation: (a) monolithic SiC,(b) MWCNTs/SiC composite, (c) SiC-coated MWCNTs/SiC composite.

10 mm 10 mm 10 mm


pullouts of MWCNTs are easily seen compared with those of the SiC-coatedMWCNTs/SiC composite. These morphological differences of the fracturedsurface indicate that SiC layers on MWCNTs should improve the weakadhesion between MWCNTs and the SiC matrix.

10.13 SEM images of the fractured surface for (a) monolithic SiC,(b) MWCNTs/SiC composite, and (c) SiC-coated MWCNTs/SiCcomposite.

(a)

(b)

(c)

1 mm

1 mm

1 mm


10.4 Conclusions

Nanometer-sized b-SiC granules are coated on diamond particles and MWCNTsuniformly by a new simple method using SiO. The SiC-coated diamondparticles and MWCNTs can act as effective reinforcements to produce newhigh-performance composites. The results can be summarized as follows.

∑ The SiC layer grows in two steps. In the first step, a thin SiC layer isformed by the direct reaction between SiO(g) and diamond or MWCNTs.In the second step, nanometer-sized SiC granules are deposited on the SiClayer by the vapor phase reaction between SiO(g) and CO(g).

∑ The oxidation resistance of diamond particles and MWCNTs is markedlyimproved by SiC coating.

∑ Dense composites of cemented carbide containing SiC-coated diamondparticles can be fabricated without conversion of diamond to graphite.The fracture toughness of the composite is double that of cemented carbidedue to the deflection and blocking effects against crack propagation bythe dispersed diamond particles.

∑ The microhardness of WC-10wt%Co increases by the dispersion of SiC-coated MWCNTs. The SiC-coated MWCNTs treated at 1150∞C to 1350∞Ccan act as nano-reinforcements for WC–Co compacts.

∑ The dispersion of SiC-coated MWCNTs increases the microhardness andfracture toughness of SiC. The SiC coating on MWCNTs at 1150∞C iseffective in improving the weak adhesion between MWCNTs and the SiCmatrix. SiC-coated MWCNT/SiC composites show elastic behavior dueto the crack-bridging effect of MWCNTs.

10.5 References

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481–486.3. Dresselhaus, M.S., Dresselhaus, G., (1997), ‘Nanotechnology in carbon materials’,

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23. Shimoo, T., Morisada, Y., Okamura, K., (2002), ‘Active-to-passive oxidation transitionfor polycarbosilane-derived silicon carbide fibers heated in Ar–O2 gas mixtures’, J.Mater. Sci., 37, 1793–1800.

24. Shimoo, T., Morisada, Y., Okamura, K., (2002), ‘Oxidation behavior of Si-M-C-Ofibers under wide range of oxygen partial pressures’, J. Mater. Sci., 37, 4361–4368.

25. Laurent, C., Peigney, A., Dumortier, O., Rousset, A., (1998), ‘Carbon nanotubes–Fe–alumina nanocomposites. Part II: Microstructure and mechanical properties ofthe hot-pressed composites’, J. Eur. Ceram. Soc., 18, 2005–2013.

26. Thostenson, E.T., Ren, Z.F., Chou, T.W., (2001), ‘Advances in the science andtechnology of carbon nanotubes and their composites: a review’, Comp. Sci. Technol.,61, 1899–1912.

27. Ma, R.Z., Wu, J., Wei, B.Q., Liang, J., Wu, D.H., (1998), ‘Processing and propertiesof carbon nanotubes–nano-SiC ceramic’, J. Mater. Sci., 33, 5243–5246.

28. Miyamoto, Y., Lin, J., Yamashita, Y., Kashiwagi, T., Yamaguchi, O., Moriguchi, H.,Ikegaya, A., (2000), ‘Reactive coating of SiC on diamond particles’, Ceramic Eng.and Sci. Proc. 21, 185–192.


29. Morisada, Y., Moriguchi, H., Tsuduki, K., Ikegaya, A., Miyamoto, Y., (2004), ‘Growthmechanism of nanometer sized SiC and oxidation resistance of SiC-coated diamondparticles,’ J. Am. Ceram. Soc., 87, 809–813.

30. Morisada, Y., Moriguchi, H., Tsuduki, K., Ikegaya, A., Miyamoto, Y., (2004), “Oxidationresistance of multiwalled carbon nanotubes coated with silicon carbide’, J. Am.Ceram. Soc., 87, 804–808.

31. Miyamoto, Y., Kashiwagi, T., Hirota, K., Yamaguchi, O., Moriguchi, H. Tsuduki, K.,Ikegaya, A., (2003), ‘Fabrication of new cemented carbide containing diamond coatedwith nanometer-sized SiC particles’, J. Am. Ceram. Soc., 86, 73–76.

32. Morisada, Y., Miyamoto, Y., (2004), ‘SiC-coated carbon nanotubes and their applicationas reinforcements for cemented carbides’, Mater. Sci. Eng. A, 381, 57–61.

33. Morisada, Y., Takaura, Y., Hirota, K., Yamaguchi, O., Miyamoto, Y., ‘Mechanicalproperties of SiC composites incorporating SiC-coated multi-walled carbon nanotubes’,J. Am. Ceram. Soc., submitted.

34. Miyata, M., Sawai, Y., Yasutomi, Y., Kanai, T., (1998), ‘Microstructure of Si3N4–SiC ceramics prepared from Si–SiO–C mixed powder’, J. Ceram. Soc. Japan, 106,815–819.

35. Fujii, K., Nakano, J., Shindo, M., (1995), ‘Evaluation of characteristic properties ofa newly developed graphite material with a SiC/C composition gradient’, Proc. 3rdInt. Symp. on Structural and Functionally Gradient Materials, ed. B. Ilschner and N.Cherradi, Lausanne, Switzerland, pp. 541–547.

36. Shimoo, T., Mizutaki, F., Ando, S., Kimura, H., (1988), ‘Mechanism of formation ofSiC by reaction of SiO with graphite and CO’, J. Japan Inst. Metals 52, 279–287.

37. Schneider, B., Guette, A., Naslain, R., Cataldi, M., Costecalde, A., 1998), ‘A theoreticaland experimental approach to the active-to-passive transition in the oxidation ofsilicon carbide’, J. Mater. Sci., 33, 535–547.

38. Narushima, T., Goto, T., Iguchi, Y., Hirai, T., (1991), ‘High-temperature active oxidationof chemically vapor-deposited silicon carbide in an Ar–O2 atmosphere’, J. Am.Ceram. Soc., 74, 2583–2586.

39. Shimoo, T., Takeuti, H., Okamura, K., (2001), ‘Thermal stability of polycarbosilane-derived silicon carbide fibers under reduced pressures’, J. Am. Ceram. Soc., 84,566–570.

40. Shimoo, T., Morisada, Y., Okamura, K., (2003), ‘Suppression of active oxidation ofpolycarbosilane-derived silicon carbide fibers by preoxidation at high oxygen pressure’,J. Am. Ceram. Soc., 86, 838–845.

41. Gargin, B.G., (1982), ‘Thermal destruction of synthesis diamond’, Advanced Materials,2, 17–20.

42. Saito, Y., Mizushima, R., Hata, K., (2002), ‘Field ion microscopy of multiwallcarbon nanotubes: observation of pentagons and cap breakage under high electricfield’, Surface Science, 499, 119–123.

43. Suzuki, H., (1986), Cemented Carbide and Sintered Hard Material, Tokyo: Maruzen.44. Omori, M., (2000), ‘Sintering, consolidation, reaction and crystal growth by the

spark plasma system (SPS)’, Mater. Sci. Eng. A, 287, 183–188.45. Takeuchi, T., Tabuchi, M., Kondoh, I., Tamari, N., Kageyama, H., (2000), ‘Synthesis

of dense lead titanate ceramics with submicrometer grains by spark plasma sintering’,J. Am. Ceram. Soc., 83, 541–544.

46. Touloukian, Y.S., Kirby, R.K., Taylor, R.E., Lee, T.Y.R., (1977), in ThermophysicalProperties of Matter, Vol. 13, Thermal Expansion. IFI/Plenum, New York, p. 19.


47. Upadhyaya, G.S., (2001), ‘Materials science of cemented carbides: an overview’,Mater. Des., 22, 483–489.

48. Faber, K.T., Evans, A.G., (1983), ‘Crack deflection process-I. Theory, and II.Experiment’, Acta Metall., 31, 565–584.

49. Moriguchi, H., Tsuzuki, K., Itozaki, H., Ikegaya, A., Hagiwara, K., Takasaki, M.,Yanase, Y., Fukuhara, T., (2001), ‘Fabrication and applications of high-toughness,highly wear-resistant diamond-and cBN-dispersed cemented carbide’, Sei TechnicalReview, 51, 121–125.

285

11.1 Introduction

While modern ceramics have unique combinations of high strength, hardness,wear and corrosion resistance, and depending on the specific system functionalproperties, their poor resistance to crack propagation can lead to catastrophicfailure under mechanical loading. As a result there has been a rich history ofattempts to improve the fracture toughness while maintaining high fracturestrengths by forming ceramic matrix composites, reinforced with secondaryphases, fibers or whiskers. One of the more recent developments in the fieldof ceramic matrix composites is the subject of nanocomposites. The termnanocomposites usually refers to a material in which sub-micron secondaryphase particles are dispersed in a polycrystalline matrix of micron grain size.

The initial focus on ceramic matrix nanocomposites was based onpolycrystalline a-Al2O3 reinforced with SiC nano-particles, which accordingto some researchers showed large increases in fracture strength (and someimprovement in toughness) as compared to monolithic alumina.1 A separateseries of nanocomposites are those based on a polycrystalline ceramic matrix,but which contain metallic nano-particles dispersed throughout the composite(a type of ceramic–metal composite, or cermet). This type of nanocompositewas first produced by the simple mixing of two oxide powders, for examplea-Al2O3 and sub-micron sized NiO, followed by sintering in a reducingatmosphere to produce sub-micron Ni particles within an alumina matrix.2

Again, cermet nanocomposites were reported to have extremely high valuesof fracture strength, and additional functional (i.e. magnetic) properties.3 Itis the goal of this chapter to explore various processing methods for theproduction of cermet nanocomposites, and to discuss in detail variations inthe microstructure which can significantly influence the final properties.

11.2 Processing

The main goal of most ceramic sintering routes is to obtain near-theoretical

11Processing and microstructural control of

metal-reinforced ceramic matrixnanocomposites

W D K A P L A N and A A V I S H A I,Technion – Israel Institute of Technology, Israel


density while controlling grain growth. This is a difficult task, since bothprocesses depend on thermally activated mass transport mechanisms whichoften occur simultaneously.4,5 For several monolithic ceramic systems, dopantshave been found which ‘enhance’ sintering, limit pore detachment fromgrain boundaries (pore occlusion), and limit grain growth, although themechanism by which such beneficial dopants work is still not clear. In thecase of a-Al2O3, MgO is the dopant of choice. Magnesium additions apparentlyincrease the solubility limit of detrimental impurities, such as Si and Ca,which form glassy phases and amorphous intergranular films,6 and/or changethe manner in which glassy phases affect grain growth.7,8 Abnormal orexaggerated grain growth is thus prevented, and a finer microstructure isobtainable after sintering to full density.

The introduction of secondary phase particles, especially metallic particles,adds a new parameter to control the evolution of the sintered microstructure,and to obtain new or refined properties. Important microstructural parametersinclude density, particle size, shape, location and content, as well as matrixgrain size which is influenced by the presence of the particles. At the sametime the introduction of metallic particles necessarily complicates the processingroute, including difficulties in slip processing, reaching homogeneousdistributions of particles in the microstructure, and very real health hazards.In the sections below, various processing routes are described, including adiscussion of these important issues.

11.2.1 Simple powder mixing

The simplest concept for the preparation of cermet nanocomposites wouldbe to use conventional mixing or milling processes, in a method analogousto that used to prepare ceramic nanocomposites, such as Al2O3 reinforcedwith SiC. The obvious problems in such a process would be the oxidation ofthe metal particles during the powder processing, prior to sintering, compoundedwith the need to deflocculate the particles. While nanometer-sized metallicparticles can be produced, the high surface area means there is a very realdanger of explosive oxidation. This effect, compounded with the relativelyhigh cost of such starting materials, limits this process to laboratory-scaleexperiments. Even laboratory-scale processes of such sorts are restricted,due to the requirement to break up soft agglomerates, which usually requiresthe additions of deflocculants. As the particle size decreases and the totalmetal surface area increases, the amount of deflocculant necessarily increases,resulting in significant amounts of long-chain molecules in the green body,which must be removed prior to or during sintering. Since the deflocculationstage is very different for metallic particles compared to the ceramic matrixparticles, there has been limited advance in this direction.

Having said this, there have been experiments to produce ceramic matrix

Processing and microstructural control 287

composites reinforced with micron-sized metallic powders, prepared by simplepowder mixing, followed by controlled atmosphere sintering. An example isthe work by Tai et al., describing alumina reinforced with nanometer-sizedCo, prepared by simple agate milling in a mixture of methanol and ethyleneglycol, followed by hot-pressing.3 Hot-pressing is required to overcome therelatively low green densities. In this study it was found that Co additionsfrom 30–50 wt% resulted in significant particle coalescence, reaching particlesizes of ~900 nm, while Co additions of up to 10 wt% resulted in Co particlesof 500 nm or less (the starting Co particle size had an average of ~30 nm).This work nicely demonstrates the main problem in processing metal-reinforcedceramic matrix nanocomposites, i.e. reaching full density while retaining thenanometer-sized length scale of the reinforcing phase. This effect is amplifiedwhen using metals with a relatively low melting point temperature, such ascopper.

11.2.2 Oxide reduction

An alternative processing method would be to utilize conventional powderprocessing of nanometer-sized oxide particles, combined with the matrixphase.2,9,10 The particles are then reduced to the metallic state during sintering,or immediately prior to sintering by using an appropriate atmosphere. Sinceboth phases are initially in the oxide state, conventional deflocculationtechniques can be combined with conventional powder processing to producegreen bodies with high densities and a homogeneous distribution of thereinforcing phase.

After production of the green bodies a reduction stage is required toreduce the oxide particles to the metallic state. The accepted approach hasbeen to sinter under reducing conditions such that the reinforcing phasereaches the metallic state during the initial stages of sintering, as the openpores in the ceramic matrix are closing. Once the ceramic matrix has reachedthe point that only closed pores exist, coarsening of the reinforcing metallicphase is limited by grain boundary diffusion kinetics through the ceramicmatrix, which is slower than surface diffusion. This of course will not preventcoalescence and coarsening of particles during the course of matrix graingrowth.

Naturally the final size of the metallic reinforcing particles is limited bythe initial size of the oxide particles. This is rather important, since there isa critical maximum particle size which can lead to degradation of the compositeproperties via thermal stress-induced cracking.2 The critical maximum particlesize was evaluated by Kolhe et al.,11 both experimentally and via finiteelement analysis. Assuming perfectly spherical particles, the difference inthermal expansion coefficients between Ni and a-Al2O3 resulted in a criticalparticle size for an isolated Ni particle of 3.0 mm. Experimental observations


for the same system showed cracking for (non-spherical shaped) particleswith a diameter of 750 nm. Thus the oxide reduction process to form metal-reinforced ceramic matrix nanocomposites still requires sub-micron ornanometer-sized oxide particles, encompassing health problems and relativelyhigh costs for the raw materials.

A relatively simple method to overcome the need for expensive andpotentially dangerous nanometer-sized oxide particles is to chemically depositthe oxide particles during the powder processing stage.12–14 This can beaccomplished quite easily by adding to water-based slips nitrates (i.e. nickelnitrate), which is calcined to the oxide state after drying. Such processesusually result in very fine oxide particle size distributions, which are reducedeither prior to or during the final sintering process. While an extra sievingand milling stage is required after calcining, water-soluble metal salts areusually inexpensive, making the process more commercially feasible.

11.2.3 Sol-gel and gel-casting

Another option increasingly being encountered in ceramic processing is theuse of sol-gels. The use of sol-gels for processing of metal–ceramic compositesintroduces a wide range of new possibilities, including the preparation ofcomplex shapes by gel-casting,15 and the option to obtain unique functionalproperties (electrical, optical, magnetic) by co-precipitation.16 This methodoffers a number of advantages, including the possibility for low temperatureprocessing, better control over homogeneity and particle dispersion, andrelatively low cost. However, this process suffers from rather low final densities,and if high-temperature sintering is involved then this usually results incoarsening of the microstructure, resulting in many cases with limitedadvantages over other processing methods.

Sol-gel processing involves the use of a hydrolysis reaction to obtain across-linked network, which results in the formation of a gel. When preparingmetal-ceramic composites, both components may be obtained in this way, oralternatively the metal reinforcement can be introduced by adding, for example,metal nitrates.17,18 The gel properties may be controlled by adjusting the pHlevel, water to metal ratio, and temperature.

The possibility to obtain a uniformly dispersed composite powder wasshown for the a-Fe–Al2O3 system where metal particles with an average sizeof 55 nm were formed in an amorphous/nano alumina matrix.18 Otherstudies attempting to obtain dense bulk composites based on the sol-gelroute using conventional pressure-assisted sintering (~1400∞C and an appliedforce of 10 MPa) resulted in a coarse microstructure.16 However, if reachingtheoretical density is not a necessary requirement, a porous ceramicmicrostructure containing nanometer-sized metal particles can be used as acatalytic material.19 Certain combinations of composite materials demand


special attention during the annealing treatments with regard to temperatureand atmosphere to avoid reduction of both species or even total loss of oneof the components.20

Gel-casting

Niihara et al.15 applied gel-casting to nanocomposite processing. Using amixture of oxide powders prepared from an alumina powder and Ni nitratedescribed above, together with methacrylamide as the monomer and N, N¢-methylenebisacrylamide as the cross-linker, they obtained a viable gel, whichwas cast and sintered to a final density of ~99%. Due to the complex objectshape, pressureless sintering was used. Niihara et al. reported a fracturestrength of ~590 ± 50 MPa for these samples, which is slightly higher thanthat of monolithic alumina. This process has the advantage of obtaining anear net-shaped object with complicated geometries while avoiding the needfor costly machining of a hard composite material.

11.2.4 Salt infiltration

A simple variant of the various methods described above is based on metalsalt infiltration into porous ceramic preforms, followed by reduction andsintering under controlled atmosphere. This method skips the more complicatedstages of calcining, secondary milling, and sieving.

The process begins with conventional ceramic powder processing to reacha green body. Partial sintering (firing) is used to induce necking at the particlecontact points, which results in a minimal level of mechanical strength,required for subsequent handling. The fired body is then infiltrated withmetal salts in a water-based solution. If the contact angle of the salt solutionis low enough (nominally under ~50∞ depending on the geometry of theparticles21), spontaneous infiltration is possible, although infiltration undervacuum is usually required to ensure complete penetration. After drying, thepreform is heated under a reducing atmosphere to form metallic particles,and sintered to full density.22

While this process is extremely simple, and various sintering modes arepossible, the amount of metal salt which can be introduced in a homogeneousmanner is dependent on the solubility limit in the liquid medium. Nitratesare extremely soluble in water, thus water-based solutions are convenient fora number of different metallic salts. However, too high a concentration canresult in large salt particles after drying, leading to micron-sized metal particles.Thus the preferred processing method is to conduct multiple infiltration–drying–reduction stages, where each stage adds ~2 wt% of metal particles tothe fired body.23 In this way various particle concentrations are possible,while the small (nanometer or sub-micron) particle size is maintained.


Furthermore, different salts can be used in the same process, allowing for theproduction of alloyed particles and/or functionally graded nanocomposites.

11.3 Microstructure

Since the very nature of nanocomposites depends on well-defined processingroutes to achieve a specific microstructure, detailed characterization of themicrostructural features is extremely important. This is particularly criticalin identifying the role of the microstructure in defining the final bulk properties.The microstructural features in nanocomposites which have been linked tobulk properties include the matrix grain size, the reinforcing particle size, itsdistribution and location (grain boundaries or occluded within the matrixgrains), segregation at the various interfaces, and residual stress fields. Figure11.1 schematically illustrates the microstructural parameters most importantto the processing and properties of metal-reinforced nanocomposites.

There are two main characterization issues which must be addressedregarding these materials. First, it is clear that conventional methods oftendo not provide the relevant information required to understand the

DMOccluded

POccluded

MTriple

PTriple

MGB

Chem

G/IGF

G/IGF

PGB

Chem

11.1 Schematic drawing illustrating the important microstructuralparameters for processing and properties of metal-reinforcednanocomposites. These include occluded pores (POccluded) and metalparticles (MOccluded), particles at grain boundaries (PGB) and triplejunctions (PTriple), segregation at grain boundaries/interfaces andintergranular films at grain boundaries/interfaces (G/IGF), as well asbulk chemistry for both the ceramic matrix and reinforcing particles(Chem).


microstructures. For example, conventional backscattered electron (BSE)images in scanning electron microscopy (SEM) simply do not have theresolution required to measure nanometer (even tens of nanometers) particlesize, for metallic particles in a non-conducting matrix. This is demonstratedin Fig. 11.2. Characterization of fracture surfaces by low-voltage SEM yields(in some cases) important morphological information regarding the locationof the reinforcing particles (triple junctions and grain boundaries) and themode of fracture (intergranular or transgranular). Figure 11.3 shows an example

2mm

11.2 A conventional SEM (backscattered electrons) micrograph of aNi–Al2O3 nanocomposite demonstrating problematic use of the BSEsignal for microstructure characterization.

Cu

500 nm

11.3 SEM (secondary electrons) micrograph of a fracture surfacefrom a Cu–Al2O3 nanocomposite. The Cu particles occupyintergranular positions and the fracture observed is mostlyintergranular.


of a fracture surface of a Cu–Al2O3 nanocomposite. In this case the locationof the metal particles could not be revealed by thermal or chemical etching.

As a result of these limitations more indirect characterization methods,such as crystal size analysis from X-ray diffraction (XRD), or methods withlimited statistical meaning but higher spatial resolution, such as transmissionelectron microscopy (TEM), must be applied. Secondly, given the enormousamount of interfacial area in nanocomposites, understanding the atomisticstructure, chemistry, and energetics (adhesion) at the metal–ceramic interfacesis a fundamental issue for nanocomposites, which to a large part has beenignored. In the following we attempt to define some of the criticalmicrostructural parameters and how they can be addressed.

11.3.1 Evolution and control

While the particle size depends initially on the method used to introduce theparticles, subsequent thermal cycles (i.e. sintering to obtain a dense composite)can lead to significant particle coarsening. This is especially so for metalparticles in a ceramic matrix, for which the sintering temperature can beclose to, if not above, the melting point of the particles. In general, when aceramic preform is sintered, the open porosity closes, and then during thefinal sintering stages the remaining closed pores are removed by vacancydiffusion through the grain boundaries. As such, the critical stage for particlecoarsening will be when open pores exist in the preform. Particle coarseningduring the final stages of sintering requires metal (cation) diffusion throughthe ceramic grain boundaries, which can occur but will be significantlyslower than surface diffusion in open pores.

While particle size is usually considered an important factor in definingthe bulk properties of the nanocomposite, particle shape can be equallyimportant, especially for mechanical properties. Normally nanometer-sizedmetal particles will be equiaxed, with a certain degree of faceting dictated bythe degree of surface energy anisotropy. However, since the particles areconfined by interfaces with the ceramic phase, the shape will depend on theinterface energy anisotropy, which is usually not known a priori, but can bedefined with a modified Wulff construction.24 For FCC metals with a limiteddegree of surface anisotropy in an alumina matrix, the interface planes areusually dictated by the facet planes of the alumina matrix.25 Similar observationshave been made regarding SiC particles in an alumina matrix.26 However,the interface energy will be affected by segregation effects, including gasspecies in the processing atmosphere. This can have a significant influenceon the shape of metal particles in a ceramic matrix, which has been demonstratedfor Cu particles in alumina due to adsorption of oxygen to the interface (seeFig. 11.4).27,28

There are several reports in the literature stating that matrix grain size


refinement, due to the existence of nanometer-sized particles at the grainboundaries, is the main reason for improved mechanical strength.14 Whilethe exact mechanism for increased strength in bending is probably morecomplicated, depending upon a combination of microstructural features, thereis no doubt that control of the matrix grain size by particle pinning is apossible advantage to nanocomposites in general. This is actually just a newlabel to a microstructural refinement mechanism which has been employedby both metallurgists and ceramicists for many years.

Grain growth kinetics depend on the driving force for grain boundarymovement, as well as the mobility of grain boundaries. During sintering (orany high-temperature annealing) the driving force for grain growth is thesimple reduction of grain boundary energy (or area).4,5 Grain boundary mobilitycan be separated into intrinsic and extrinsic mobility. Intrinsic mobility willdepend on the exact mechanism for atom transfer from one ‘side’ of theboundary to the other. Extrinsic mobility depends upon a variety of factors,which include solute drag due to grain boundaries (this is a result of segregationmoving with the boundary), impurity (dopant) segregation, mobility due tothe presence of a liquid phase at the boundary which increases local diffusionrates, and grain boundary pinning by secondary phases (known as Zener drag).29

Grain boundary pinning by secondary particles is naturally a point ofinterest for processing nanocomposites. Grain boundary pinning can bedescribed as a drag force on the grain boundary, which negates the drivingforce for grain growth, and was elegantly described by Hsueh et al.30 via thedihedral angle y formed at the triple junction of a particle with radius r at agrain boundary with energy gGB:

F = p rgGB(17.9 – 6.2y) (11.1)

Assuming the particle remains at the grain boundary, and is not occluded bythe growing crystal, the larger the particle and the smaller the dihedral angle,the larger the drag force negating grain growth. For nanocomposites it is theaccumulated force of many small particles acting at each grain boundary

Cu

500 nm

11.4 TEM micrograph of a Cu–Al2O3 nanocomposite. The Cu particlesare elongated and located at grain boundaries and triple junctions.


which defines the drag force, compensating for the small particle radius. Thedihedral angle is a measure of the interface energy, thus elongated particles(as shown in Fig. 11.4) result in a high drag force and more effectively limitmatrix grain growth.

11.3.2 Interfaces

Due to small particle size, the amount of metal–ceramic interface area innanocomposites can be quite significant, even for metal particle additions ofless than 10 wt%. As such, the nature of the metal–ceramic interfaces caninfluence the bulk properties to extents not normally achieved in conventionalmetal–ceramic composites. Relevant parameters include the interface energy,which defines the thermodynamic work of adhesion and has a significantand non-linear connection with interface toughness.31 The interface energywill in turn be influenced by segregation of constituent atoms and impurities/dopants, including elements in the processing atmosphere.

Basic studies of metal–ceramic interface thermodynamics have clearlydemonstrated the importance of segregation on adhesion. An example pertinentto Ni-reinforced alumina matrix nanocomposites is the influence of Alsegregation on the Ni–alumina interface energy, where additions of only afew atomic percent of Al to Ni can significantly increase the thermodynamicwork of adhesion.32,33 Pre-doping the Ni particles is not necessary, sincesome dissolution of alumina is expected as the particle and alumina grainsadjust their shape to reduce the total interface area through the formation oflenticular-shaped particles at grain boundaries.34 While this is an importantissue in basic surface or interface science, it has yet to be investigated orapplied to nanocomposites.

One interesting example of possible modification of interfaces in metal-reinforced nanocomposites is the presence of intergranular films. Highresolution transmission electron microscopy (HRTEM) of grain boundariesand interfaces in such materials revealed the existence of thin, apparentlyamorphous films, with a constant thickness in the range of 1–2 nm.25 Theexistence of thin intergranular films at grain boundaries in ceramics is awell-established phenomenon, and has been characterized in detail at grainboundaries in alumina, silicon nitride, and silicon carbide.7,35,36 The formationof the film is usually a result of sintering additives or impurities, which formamorphous phases commonly found at triple junctions.

Clarke and co-workers developed a model to calculate the thickness ofthe amorphous film observed in polycrystalline ceramics.37,38 The model isbased on a force balance between an attractive van der Waals dispersionforce that acts across the grain boundaries, any capillary forces present, andrepulsive disjoining forces (such as steric forces and electrical double-layerforces) in the amorphous film.37,38 The repulsive steric force is based on the


assumption that the first monolayer of molecules of the glassy phase exhibitsa heteroepitaxial arrangement with both grain surfaces and that this orderingextends over a certain distance within the amorphous film. The electricaldouble-layer forces occur if the interface between the grains and the film ischarged by cations. While the van der Waals and capillary forces act to bringthe grains closer together, the disjoining forces widen the film. The estimatedvalues of the amorphous film thickness, which is dependent on the dielectricproperties of the film and the grains, are in excellent agreement with theexperimentally observed film width.37,39,40 Due to the balance of forces leadingto a defined film thickness, such films are often termed equilibrium amorphousfilms.

Recently, it was shown that equilibrium amorphous films also exist atinterfaces between metals (Ni and Cu) and alumina, via model experimentsbased on nanocomposites. Cu–Al2O3 and Ni–Al2O3 nanocomposites wereprepared using high-purity alumina powder to which predetermined amountsof Ca and Si dopants were added.25

Detailed HRTEM characterization of the specimens showed that all metal–ceramic interfaces in the two different nanocomposites had thin (~1 nmthick) amorphous films (see Fig. 11.5). In addition, occluded particles werefound inside the alumina grains which also had thin amorphous films at theirinterfaces with alumina. Analytical microscopy showed the films to containCa, Si, and Al.41 Hamaker coefficients were calculated for metal–ceramicinterfaces in the presence of a SiO2-based film, which indicated that a strongerattractive force is expected for intergranular films at metal–alumina interfaces,

11.5 HRTEM micrograph of a Ni–Al2O3 interface taken from a Ni–Al2O3 glass-doped nanocomposite. A thin (~1 nm) amorphous filmexists at the metal–ceramic interface, extending from a glass-pocketat the triple grain boundary junction.

Ni

Al2O3

5 nm

(111)

(0001)


relative to alumina grain boundaries, correlating to the experimentally measuredfilm thickness.40

The ability to modify the metal–ceramic interface in nanocomposites bythe formation of intergranular films holds exciting prospects. From athermodynamic point of view, the existence of a film at equilibrium indicatesa lower interface energy than an interface without a film. This indicates thepotential to increase the adhesion of interfaces, although experimentalinvestigations are required to fully evaluate this effect. However, the promotionof particle occlusion due to the presence of the films has been shown,28 andthis means that a new method to modify and control the microstructuralevolution of nanocomposites is available, as discussed in the next section.

11.3.3 Particle occlusion

As mentioned in Section 11.3.1, the particles in nanocomposites induce adrag force on the matrix grain boundaries, which in some cases can reducegrain boundary migration rates, resulting in a finer matrix microstructure.However, if the grain boundary migration rate is faster than that of theparticles, particle detachment can occur, resulting in the particle being occludedwithin the growing matrix grain. Detachment (or occlusion) depends on therelative interface mobility for a given driving force.

In the nanocomposites mentioned in the previous section, occluded metalparticles were not found in nanocomposites without the glassy phase.28

However, when the nanocomposites were doped with glass-forming elements(Ca and Si), intergranular films were detected and a significant amount ofmetal particles was found to be occluded within the alumina grains (see Fig.11.6). The occluded particles in the Cu–alumina and Ni–alumina compositeshad an average size of 250 ± 20 nm and 260 ± 90 nm respectively, while theparticles located at alumina grain boundaries had an average size of 1400 ±300 nm and 850 ± 350 nm, respectively. These measurements were performedusing TEM in order to differentiate between particles found at alumina grainboundaries from the occluded particles.

A similar type of behavior is seen in other studies. In composites preparedby Oh et al. using high-purity alumina, only a few occluded particles wereobserved.14,42 However, in a study by Chen and Tuan,43 a large number of Niparticles were reported to be occluded (~20–30%). Chen and Tuan observeda similar difference in the distribution of the reinforcing particle size betweenthe occluded particles (limited to ~100 nm) and the particles found at grainboundaries and triple junctions. Based on the findings from the samplespurposely doped with glass and containing intergranular films, resulting in70–80% of the particles occluded in the alumina grains, the question israised whether the samples prepared by Chen and Tuan43 contained smallamounts of glass-forming impurities.


Two important points should be noted regarding metal particle occlusion:the fact that glass doping promotes particle occlusion, and the distinct sizedistribution of the occluded particles relative to particles found at grainboundaries. Unless a special orientation relationship exists between the matrixand an occluded particle, the occlusion process should be energeticallyunfavorable, since it will increase the surface energy of the system. Theocclusion of particles (and pores) is therefore a kinetic result. For the particle(or pore) to stay attached to a grain boundary, its velocity should be the sameas that of the grain boundary.44 The velocity (Vi) may be described as theproduct of two parameters – the driving force (Fi) and the mobility (Mi):

Vi = Fi · Mi (11.2)

The driving force for grain growth (Fb) is the reduction in the total internalsurface (grain boundary) energy. This can be expressed by the grain boundaryenergy (gGB) and average grain size ( G ):44

FG

bGB

3ª g(11.3)

For a given grain boundary energy, the smaller the grain size the larger thedriving force for grain growth. Due to the fact that a stable intergranular filmforms during the initial stages of sintering, it may be assumed to a firstapproximation that the mobility would not change drastically during thesintering process. This would mean that the velocity of the grain boundarieswill decrease proportional to 1/ G during the sintering.

The drag force for a particle (or pore) was defined in equation (11.1). The

11.6 Bright field TEM micrograph showing the metal particlemorphology in a Cu–Al2O3 glass-doped nanocomposite.

Occluded Ni

Occluded pores

500 nm


mobility (Mp) of a particle/pore can be expressed in a similar way if theparticle is assumed to have a dominant interfacial diffusion mechanism:29,44

MDkT r

MDkT rp(s)

s s4 p (I)

I I4 = , =

dp

dp

W W(11.4)

where Ds and DI are the surface and interface diffusion coefficients respectively,ds is the boundary core width, W is the total ionic volume divided by thenumber of slow diffusing ions, k is Boltzmann’s constant, and T is the absolutetemperature. For metal-ceramic interfaces, diffusion tends to be faster thanfor ceramic surfaces.45 However, in some cases volume diffusion dominatesover interfacial diffusion, as was shown by Saiz et al.45 and Monchoux andRabkin.46 In this case mobility is a function of 1/r3:29,46

MD ckT rp(V)

V V3 =

34

Wp

(11.5)

where cV is the solubility of the diffusing atoms in the particle. Whicheverdiffusion mechanism is dominant, the velocity of small particles is expectedto be higher than that of larger ones, since the velocity will depend on 1/r3

for surface/interface diffusion while for volume diffusion it will depend on1/r2. This means the observed occluded particle size distribution (i.e. smallparticles), and larger particles found at the alumina grain boundaries, cannotbe controlled by particle velocity dependence on particle size.

Both the matrix grains and metal particles coarsen during sintering. Duringthe initial stages of sintering, due to the presence of the thin liquid film at thealumina grain boundaries, the mobility of the boundaries is increased, and atthe same time the driving force for grain growth is high since the aluminastarts from a fine 0.3 mm size powder. This results in a high velocity of thegrain boundaries, and relatively small occluded particle size.

As the sintering process advances, the metal particles coarsen, whichreduces their velocity and should result in their occlusion. However,simultaneous grain growth of the alumina grains reduces the driving forcefor grain growth and therefore the grain boundary velocity. This enables themetal particles to advance together with the moving alumina boundaries andresults in the observed particle size distribution.

This kinetically dependent mechanism provides a means to develop ananocomposite microstructure with the particles (or the majority of the particles)occluded within the matrix grains. On the other hand, occlusion can beprevented, for the most part, if glass-forming impurity elements are notintroduced into the material during the processing stage. As we will see inthe next section, the position of the particles (i.e. occluded or at grainboundaries) can influence the microstructurally dependent properties ofnanocomposites.


11.3.4 Residual stresses and fracture strength

The measurement of residual stresses is usually associated with the analysisof mechanical properties, and not microstructure. However, residual stressfields in nanocomposites depend directly on microstructural parameters (particlesize, position and spacing), as well as bulk material properties, such asdifferences in the coefficient of thermal expansion.

The residual stresses in a composite can be of two types. The first is oftencalled a macrostress, resulting from, for example, cutting, grinding or polishing.Macrostresses will result in changes in atomic spacing across a volume ofmaterial, and can be measured by the shifts (D2q) of XRD reflections. Incomposites microstresses are associated with thermoelastic mismatch betweendifferent phases or between anisotropic grains. Microstresses areinhomogeneous in nature, and will fluctuate from point to point in the material.While the average value of microstresses, across a given volume of material,will also contribute to XRD reflection shifts, fluctuations in the amplitude ofthe microstresses causes broadening of XRD reflections.47

Any nanocomposite material will have some degree of residual stress,due to the difference in thermal expansion coefficients between the particleand matrix phases. The paradigm for stress analysis in nanocomposites is theSiC-reinforced alumina system, where the smaller coefficient of thermalexpansion of the SiC particles, relative to alumina, induces in the surroundingalumina a compressive radial stress and a tangential tensile hoop stress.48

Due to the anisotropic thermal expansion of alumina, ~100 MPa tensilestress fields are expected to exist along alumina grain boundaries.49 Thethermally induced compressive stresses due to SiC particles adjacent to aluminagrain boundaries may account for the reduced pullout of grains during polishingwhich has been observed for SiC-reinforced alumina nanocomposites.50

Furthermore, tensile hoop stresses due to occluded SiC particles would beexpected to contribute to the transition from intergranular to transgranularfracture. Levin et al. explored the microstructurally dependent residual stressfields by analyzing their distribution and amplitude as a function of particlesize and distribution (or content).51 The contribution of fluctuating microstressfields to changes in the material fracture toughness showed that an increasein fracture toughness is expected only for relatively small particle contents(3.5–5 wt% SiC), which can be optimized by reducing the particle size forthe same volume fraction of occluded particles.

For most metal-reinforced nanocomposites the thermal expansion coefficientof the metal phase will be larger than that of the matrix, reversing the expectedstress fields compared to SiC-reinforced alumina. Thus while the tensileradial stresses surrounding occluded particles may induce transgranularcracking, the compressive hoop stresses may inhibit crack propagation if theparticles are located at grain boundaries. Macrostresses in sub-micron Ni


particles in a sintered alumina matrix have been measured using XRDtechniques.23 The level of stress can be quite significant, even consideringthat the yield strength of bulk Ni at high temperatures is significantly reducedcompared to room temperature (i.e. less than 10% of its room temperaturevalue at ~1000∞C).52 This is no doubt due to the limited plasticity in confinednanometer-sized particles.

11.4 Properties

The initial interest in ceramic matrix nanocomposites arose from reports byNiihara and co-workers indicating enhanced mechanical properties due tothe presence of ceramic (SiC) particles.53 With the development of variousprocessing routes to introduce nanometer-sized metal particles in a ceramicmatrix, variations in functional (i.e. magnetic) properties are possible. In thefollowing we briefly review the microstructurally dependent properties, withemphasis on the possible mechanisms leading to improved properties andusing SiC-reinforced alumina as a point of comparison.


Particle-reinforced ceramic matrix nanocomposites gained attention due totheir enhanced mechanical properties, demonstrated by Niihara and co-workers.53 The initial focus was on ceramic particle (SiC)-reinforced alumina,with matrix grain size refinement due to grain boundary pinning being theinitial explanation for the high fracture strengths. With the advent of metal-reinforced nanocomposites, matrix grain size refinement was again used toexplain the high fracture strengths (see Fig. 11.7).12

Upon examining the residual stress in SiC-reinforced alumina, it appearsthat thermally induced compressive stresses due to occluded particles inducetransgranular cracking and reduce crack propagation along the matrix grainboundaries.1 According to this model, only occluded particles, within a specificparticle concentration limit,51 contribute to the two combined strengtheningmechanisms, and particles located at grain boundaries will reduce the fracturestrength of the nanocomposite. This may be the reason for alternative processingroutes leading to fracture strengths lower than the values reported by Niiharaand co-workers, i.e. nanometer-sized particle occlusion was not achieved ora significant number of particles remained at the grain boundaries.

In metal-reinforced ceramic matrix nanocomposites, matrix grain refinementhas also been demonstrated by Niihara and co-workers.12 At the same time,as mentioned in the previous section, significant residual stress fields havebeen measured for Ni-reinforced alumina nanocomposites,23 although withreversed compression–tension fields compared to SiC-reinforced alumina.Thus for metal particles below the critical size for intrinsic cracking11 and


located at grain boundaries, compressive hoop stresses exerted by the metalparticles on the ceramic grain boundaries are expected to play some role inreducing crack propagation.22 While occluded metal particles will inducelocal tensile radial stresses in the alumina grains, it has yet to be directlyshown that this mechanism results in a change in mode to transgranularcracking.

It is most unlikely that matrix grain refinement alone can explain the highfracture strengths that have been reported. Using the reported maximumfracture strengths (s = 1100 MPa)12 and the nominal fracture toughness ofalumina (KIc = 3.5 MPa.m1/2), the critical flaw size, (c), can be estimatedfrom

cKY

= Ic2

sÊË

ˆ¯ (11.6)

to be of the order of 10 mm. This is an order of magnitude larger than thematrix grain sizes measured for the same nanocomposites. So while matrixgrain size reduction may contribute to increases in fracture strength, it canonly be part of the series of mechanisms leading to a reduced effective flawsize.

Some reports have cited bridging across the relatively ductile metal particlesas mechanisms which contribute to the increase in fracture strength.54 Otherreports demonstrate cracks propagating at the metal–ceramic interfaces,

0 5 10 15 20Ni (vol%)

1200

1100

1000

900

800

700

600

Frac

ture

str

eng

th (

MP

a)

11.7 Effect of volume fraction of Ni particles on the fracture strengthof Ni–Al2O3 nanocomposites, produced by the reduction of NiOparticles (■), and the reduction of Ni-nitrate (∑). Reproduced fromSekino et al.,12 with permission from the Journal of the AmericanCeramic Society.


precluding bridging effects.12 It is almost certain that metal particles in thesize range from nanometers to a few hundred nanometers do not contributeto bridging effects and have extremely limited plasticity. It is far more likelythat energy dissipation mechanisms of advancing cracks are defined by themetal–ceramic interfaces and the stress fields around the particles.

11.4.2 Functional properties

A thorough review of the various functional properties obtained using metal–ceramic composites is beyond the scope of this work. However, a few casesand general trends will be briefly considered. Obtaining nanometer-sizedmetal particles dispersed in a ceramic matrix is attractive for a number ofreasons. The ceramic matrix provides protection against corrosion or high-temperature oxidation of the normally oxidation-susceptible metal particles.Ferromagnetic particles in this size range can form single magnetic domains.Combining the latter with a conducting or semiconducting matrix which isstable at high temperatures can form a system with giant-magnetoresistanceproperties.23,55

The ratio of surface area to bulk volume of the reinforcing particles canhave important implications on optical properties, where the contribution ofsurface states can result in unique properties.56,57 These surface states causeshifts in the plasmon absorption frequencies and can be manipulated by useof different combinations of metals and ceramics.56 Another possibility dueto the high surface area of the metal particles is catalysis applications, providedthe ceramic matrix contains open pores.19

The key to most of the functional properties reported is a fine microstructureof the metal particles (i.e. in the nanometer scale) which is uniformly dispersedwithin a ceramic matrix. In some cases the particle size needed is in therange of a few nanometers in order to enhance the surface properties, whilein other cases optimization is needed between the demand for single domainparticles while minimizing unwanted surface states.

11.4.3 Oxidation resistance

Due to the potential high-temperature application of nanocomposites, aswell as the fact that metal-reinforced ceramic nanocomposites combine metaland non-metal phases in equilibrium, it is important to understand the oxidationresistance of such materials. Using the Ni–alumina system as an example,and following Sekino et al.,12 the partial pressure of oxygen required toprevent the formation of nickel spinel (NiAl2O4) from a two-phase mixtureof Ni and Al2O3 can be described as:58,59

PTO2 (atm) = exp

–58240 480 + 17.94

±ÊË

ˆ¯ (11.7)


Figure 11.8 graphically demonstrates this relationship, and shows that forconventional sintering temperatures (1400–1600∞C), a reasonable partialpressure of oxygen (10–8–10–6 atm) is required to prevent the reaction

Ni + Al O + 12

O NiAl O2 3 2 2 4Æ (11.8)

Such levels of oxygen are fairly simple to maintain during processing byusing a slightly reducing atmosphere, and the cooling rates employed duringprocessing are apparently fast enough to prevent spinel formation.

However, during long exposures to medium-temperature operatingconditions, e.g. 1000∞C, spinel formation is certainly expected. Wang et al.60

demonstrated this for the Ni–alumina system, showing the diffusion of Niatoms to the free surface of the nanocomposite, followed by the formation ofa nickel spinel surface coating which then limits the kinetics of subsequentoxidation. In this case the formation of a spinel surface layer may be beneficialto mechanical properties, since the reaction results in a volume increase, andthe formation of compressive residual stresses. An analogous behavior wasreported for ceramic particle nanocomposites, where oxidation of SiC particlesresults in an increase in volume and compressive residual stresses.61

In an opposite trend, the addition of ZrO2 particles to Ni-reinforced alumina

11.8 Graph showing the partial pressure of oxygen required toprevent the formation of nickel spinel (NiAl2O4) from a two-phasemixture of Ni and Al2O3 as a function of temperature.

Ni +Al2O3

NiAl2O4 + Al2O3

800 1000 1200 1400 1600 1800Temperature (∞C)

10–5

10–7

10–9

10–11

10–13

10–15

10–17

10–19

10–21

10–23

10–25

P(O

2) (

atm

)


matrix nanocomposites degraded the oxidation resistance of thenanocomposite.62 Above the critical ZrO2 particle concentration for percolation,the ZrO2 phase provides a rapid route for oxygen diffusion, which enhancesoxidation of the composite.

11.5 Future trends

It is very clear from the numerous studies on metal-reinforced ceramic matrixnanocomposites that controlling the microstructural features is critical tocontrolling and/or achieving specific properties. At the same time, ourunderstanding of the mechanisms controlling microstructural evolution insuch complicated systems is limited. Basic studies are required to understandthe mechanisms, and specifically the role of dopants (including gas phases)and especially metal cations on sintering and grain boundary migration rates.In a similar manner, correlating interface energy with fracture energy atmetal–ceramic interfaces is critical to the design of optimized mechanicalproperties, while interface structure and chemistry is important for functionalproperties.

Alternative processing methods also offer the potential to control themicrostructure and final properties of nanocomposites. Both self-propagatinghigh-temperature sintering and spark plasma sintering offer means to obtainmetastable yet dense nanocomposites. Subsequent heat treatments can thenbe used to approach equilibrium microstructures, where the properties willbe a function of the heat treatment temperature and time. In this way avariety of microstructures, and thus variations of the composite properties,can become available.

Additions of more than one type of metal particle, intermetallic phases, orgraded particle concentrations offer a rich field for research into potentialfunctional properties. Based on the processes discussed in this chapter, it isfairly simple to introduce different types of metallic particles into the ceramicmatrix, which depending on the respective phase diagram would either remainas separate phases, form solid solutions, or form intermetallic compounds.Varying the partial pressure of gas components in the sintering atmospherefurther expands the number of degrees of freedom to form different nanometer-sized phases. Finally, as demonstrated, intergranular films can form both atthe matrix grain boundaries and between the matrix grains and the reinforcingparticles, which can alter the processing kinetics and final properties of themetal-reinforced ceramic matrix nanocomposite.

11.6 References

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24. Gabrisch, H., Dahmen, U. and Johnson, E., ‘In-situ observations of the interactionof liquid lead inclusions with grain boundaries in alumina’, Microscopy andMicroanalysis, 1998 4 286–293.

25. Avishai, A., Scheu, C. and Kaplan, W.D., ‘Amorphous films at metal–ceramicinterfaces’, Zeitschrift für Metallkunde, 2003 94 272–276.

26. Kaplan, W.D., Levin, I. and Brandon, D.G., ‘Significance of faceting on SiCnanoparticles in alumina’, Mater. Sci. Forum., 1996 207–209 733–736.

27. Ghetta, V., Fouletier, J. and Chatain, D., ‘Oxygen adsorption isotherms at the surfacesof liquid Cu and Au–Cu alloys and their interfaces with Al2O3 detected by wettingexperiments’, Acta. Mater., 1996 44(5) 1927–1936.

28. Avishai, A. and Kaplan, W.D., ‘Intergranular films in metal–ceramic composites andthe promotion of metal particle occlusion’, Zeitschrift für Metallkunde, 2004 95266–270.

29. Gottstein, G. and Shvindlerman, L.S., ‘Theory of grain boundary motion in thepresence of mobile particles’, Acta. Mater., 1993 41(11) 3267–3275.

30. Hsueh, C.H., Evans, A.G. and Coble, R.L., ‘Microstructure development duringfinal/intermediate stage sintering – I. Pore/grain boundary separation’, Acta. Metall.,1982 30 1269–1279.

31. Lipkin, D.M., Clarke, D.R. and Evans, A.G., ‘Effect of interfacial carbon on adhesionand toughness of gold–sapphire interfaces’, Acta. Mater., 1998 46(13) 4835–4850.

32. Levi, G., Scheu, C. and Kaplan, W.D., ‘Segregation of aluminium at nickel–sapphireinterfaces’, Inter. Sci., 2001 9 213–220.

33. Levi, G., Clarke, D.R. and Kaplan, W.D., ‘Free surface and interface thermodynamicsof liquid nickel in contact with alumina’, Inter. Sci., 2004 12(1) 73–83.

34. Saiz, E., Tomsia, A.P. and Cannon, R.M., ‘Ridging effects on wetting and spreadingof liquids on solids’, Acta. Mater., 1998 46(7) 2349–2361.

35. Kleebe, H.J., Hoffmann, M.J. and Rühle, M., ‘Influence of secondary phase chemistryon grain boundary film thickness in silicon nitride’, Zeitschrift für Metallkunde,1992 83(8) 610–617.

36. Moberlychan, W.J., Cao, J.J, and De, Jonghe, L.C., ‘The roles of amorphous grainboundaries and the beta–alpha transformation in toughening SiC’, Acta. Mater.,1998 46(5) 1625–1635.

37. Clarke, D.R., ‘On the equilibrium thickness of intergranular glass phases in ceramicmaterials’, J. Am. Ceram. Soc., 1987 70(1) 15–22.

38. Clarke, D.R., Shaw, T.M., Philipse, A.P. and Horn, R.G., ‘Possible electrical double-layer contribution to the equilibrium thickness of intergranular glass films inpolycrystalline ceramics’, J. Am. Ceram. Soc., 1993 76(5) 1201–1204.

39. Choi, H.J., Kim, G.H., Lee, J.G. and Kim, Y.W., ‘Refined continuum model on thebehavior of intergranular films in nitride ceramics’, J. Am. Ceram. Soc., 2000 83(11)2821–2827.


40. Avishai, A. and Kaplan, W.D., ‘Intergranular films at metal–ceramic interfaces: PartII – Calculation of Hamaker coefficients’, Acta. Mater., 2005 53(5) 1571–1581.

41. Avishai, A., Scheu, C. and Kaplan, W.D., ‘Intergranular films at metal–ceramicinterfaces: Part I – Interface structure and chemistry’, Acta. Mater., 2005 53(5)1559–1569.

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43. Chen, R.Z. and Tuan, W.H., ‘Pressureless sintering of Al2O3/Ni nanocomposites’, J.Eur. Ceram. Soc., 1999 19 463–468.

44. Powers, J.D. and Glaeser, A.M., ‘Grain boundary migration in ceramics’, Inter. Sci.,(1998) 6(1–2) 23–39.

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46. Monchoux, J.P. and Rabkin, E., ‘Microstructure evolution and interfacial propertiesin the Fe–Pb system’, Acta. Mater., 2002 50 3159–3174.

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48. Levin, I., Kaplan, W.D., Brandon, D.G. and Wieder., T., ‘Residual stresses in alumina–SiC nanocomposites’, Acta. Mater., 1994 42(4) 1147–1154.

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51. Levin, I., Kaplan, W.D., Brandon, D.G. and Layyous, A.A., ‘Effect of SiCsubmicrometer particle size and content on fracture toughness of alumina–SiCnanocomposites’, J. Am. Ceram. Soc., 1995 78(1) 254–256.

52. Rosenberg, S.J., ‘Nickel and its alloys’, National Bureau of Standards Monograph,1968 106 38–39.

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309

12.1 Introduction

Carbon nanotubes (CNTs) have recently emerged in the research world asone of the most promising nanomaterials. Both their characteristics and theirproperties lead one to think that their incorporation in materials could lead tonew nanocomposites with new or enhanced physical or mechanical properties.In the present chapter, the results of researches which have been recentlydeveloped to prepare and characterize novel CNT-ceramic composites arereviewed. Firstly, in Section 12.2, the different structural forms of CNTs aredescribed and the most common synthesis methods are presented as well asthe physical and mechanical properties. In Section 12.3, the different processeswhich have been retained to obtain homogeneous dispersions of CNTs inceramic powders are explained and the methods used to densify the compositesare compared. In Section 12.4, the mechanical properties, and most particularlythe results obtained for fracture toughness, are discussed, the effects of CNTaddition on the electrical conductivity of insulating or semi-conducting ceramicsare described, and works reporting the thermal conductivity of these compositematerials are presented. Then, in the light of these very recent results, thekey problems which require further researches will be discussed.

12.2 Structure, synthesis and properties of carbon

nanotubes

Five years after the discovery of fullerenes, Iijima reported in 19911 a novelform of organized carbon which consists of hollow cylindrical structures, afew nanometers in diameter and some micrometers long. Although hollowcarbon nanofibers had been prepared for several decades, their walls hadnever been resolved by High-Resolution Transmission Electron Microscopy(HRTEM). These HRTEM images allowed Iijima to conclude that the wallsof the so-called multi-walled carbon nanotubes (MWCNTs) are made up ofseveral concentric cylinders, each being formed by a graphene sheet rolled

12Carbon nanotubes-ceramic composites

A P E I G N E Y and C H L A U R E N T,CIRIMAT, Université Paul-Sabatier, France


onto itself so that the open edges match exactly. Half fullerenes or morecomplex structures that include pentagons close the tips of each cylinder. In1993, the Iijima and Bethune groups reported simultaneously2,3 the synthesisof CNTs composed of only one wall, the so-called single-walled carbonnanotubes (SWCNTs).

For each cylindrical wall, the hexagons can present a lot of differentorientations related to the tube axis, each giving a particular structure (Fig.12.1).4 Each orientation is represented by a helical vector

rCh which is deduced

from the director vectors ( and )1 2r ra a of the graphene sheet by using a pair

of integers (n, m): r r rC na mah = + 1 2 (Fig. 12.2). The limiting cases are referred

to as zig-zag (n, 0) and armchair (n, n) SWCNTs (Figs 12.1(b) and 12.1(a)respectively), with a helical (or chiral) angle q equal to 0∞ and 30∞, respectively.Other SWCNTs have a helical (often termed chiral) structure (Fig. 12.1(c)).It will be seen later that the electronic properties of CNTs greatly depend ontheir helicity. The most common defects in the wall consist of pentagons orheptagons which can induce elbows on a tube or allow junctions betweentubes of different structures. It is important to note that most SWCNTs aregenerally found in ropes (or bundles) of several tens of SWCNTs, arrangedin a triangular lattice. The intertube distances are around 0.34 nm. In MWCNTs,the measured interlayer distance (0.34–0.39 nm) is close to that measuredbetween graphene sheets in graphite and no particular correlation appears

…

…

…

(a)

(b)

(c)

12.1 Three examples of particular structures of SWCNTs, dependingon the orientation of the hexagons related to the tube axis. (a)armchair-type tube (q = 30∞), (b) zigzag type tube (q – 0∞), and chiraltube (0 < q < 30∞). Reprint from Carbon, vol. 33, No. 7, DresselhausM.S., Dresselhaus G., Saito R., Physics of carbon nanotubes, pages883–891, Copyright (1995) with permission from Elsevier.

Carbon nanotubes-ceramic composites 311

between the helicity of concentric layers. CNTs made up of two walls, theso-called double-walled nanotubes (DWCNTs), are of great interest, particularlyfor composite applications, because the outer wall provides the interfacewith the matrix and can be functionalized to control interface, whereas theinner wall is protected and thus retains its intrinsic properties.

Three main methods are used to synthesize CNTs: arc-discharge evaporationof graphite electrode, laser sublimation of graphite rods, and catalytic chemicalvapor deposition (CCVD).5 The first method produces short and well-crystallized CNTs, either MWCNTs (without a catalyst) or SWCNTs (witha catalyst) with a narrow diameter distribution (around 1.2–1.4 nm). TheSWCNTs are generally arranged in ropes. The main problem is the fairly lowpurity of the samples, which requires oxidative treatments resulting in somedamage to the CNTs. The laser method produces also ropes of SWCNTshaving a narrow diameter distribution but with a much higher purity. Theseropes, however, have a larger diameter (several tens of nanometers) thanthose of the arc-CNTs and can be made up of hundreds of SWCNTs. CCVDroutes derive from the methods used for decades for the synthesis of carbonnanofibers (CNF). The catalytic decomposition, at a high temperature (600–1100∞C), of a carbonaceous gas (generally a hydrocarbon or carbon monoxide)on nanoparticles (mainly Fe, Co or Ni) produces the formation of CNTs.

Generally, the mechanism involved in the formation of a CNT is similarto that of a CNF, i.e. only one CNT from one nanoparticle. Dependingmainly on the diameter of the active catalytic particle, and also on the reaction

B

T

O

B¢

y

x

q

A

Ch

a1

a2

12.2 Helical (or chiral) vector

rCh defined from the director vectors

( ) and ( )1 2r ra a of the graphene sheet by using a pair of integers (n, m) :

r r rC a ah = n + m and1 2 chiral angle q. Reprint from Carbon, vol. 33, No. 7,Dresselhaus M.S., Dresselhaus G., Saito R., Physics of carbonnanotubes, pages 883–891, Copyright (1995) with permission fromElsevier.


conditions, either a MWCNT, a SWCNT or a DWCNT can be formed. Thus,particular catalytic materials or methods are needed to obtain a sufficientselectivity, particularly to produce SWCNTs or DWCNTs, because it isnecessary to limit the size of the catalytic particle to a few nanometers(<3 nm), at a high temperature.

The SWCNTs or DWCNTs generally have a distribution of diameterslarger than with the previous methods. They are individual or in small ropes,and can be up to 100 mm long. MWCNTs produced by CCVD, particularlythose with many walls, often achieve less crystallization than those producedby arc or laser methods. They are generally very long (hundred of micrometers),and the presence of defects can induce important irregularities in the walls.Some are compartmented and sometimes have a bamboo shape. In this case,the designation of these carbon filaments as MWCNTs or as CNFs is debatable.CCVD methods are the most promising because they are economical andcan be scaled up. Several companies currently sell MWCNT or SWCNTsamples produced by such methods. But, before attempting to use the CNTsto prepare a composite, it is interesting to investigate the characteristics ofthe samples because each method produces CNTs with their own particularitiesand purities. Furthermore, particular attention must be paid to the eventualdamage to CNTs, or to the functionalization, which could have been inducedby purification treatments.

CNTs are unique one-dimensional hollow objects with a very high aspectratio (100–100 000) and a very high specific surface area (1300 m2/g forclosed SWCNTs6). Their ends are closed and they have a low reactivity, butoxidative treatment is efficient in opening the ends and getting functionalgroups on the walls. They combine a high Young modulus (>1 TPa) with ahigh tensile strength (mean value close to 30 GPa)7 and great resilience dueto the capacity to form kinks reversibly. These mechanical properties arerelated to well-crystallized individual CNTs, and are less for most MWCNTsproduced by CCVD. Otherwise, inside a rope, the SWCNTs are weaklylinked and allow easy sliding, which is detrimental to the tensile strength.The thermal conductivity of CNTs is very high (2000–6000 W.m–1K–1). Theelectronic properties of CNTs have been reviewed by Ahlskog et al.8 Themost important characteristic to note for use of the CNTs in composites isthat either a metallic or a semiconducting behavior may be observed, dependingon the diameter and helicity. The electrical conductivity of a metallic CNTcould reach 10 000 S.cm–1 and that of a semiconducting CNT is in the range0.1–100 S.cm–1. But no method has currently been found to produce CNTsamples containing only metallic or only semiconducting CNTs. It has onlybeen reported that some separation methods tend to discriminate SWCNTsas a function of their helicity, but only at a microscopic scale.


12.3 Preparation of CNT-ceramic composites

After getting a sufficient quantity of CNTs and collecting their characteristics,the first difficulty in preparing dense CNT-ceramic composite materials is toobtain a composite powder in which the CNTs are well distributed, withoutforming CNT aggregates. Mixing by co-milling of CNT and ceramic powders,generally in wet media, has been used by several authors. The better efficiencyof the synthesis of CNT in situ within the ceramic powders or, symmetrically,the synthesis of the ceramic in situ around the CNT, has been demonstrated.The second difficulty is to achieve good densification of the material. Bothhot-pressing (HP) and spark plasma sintering (SPS) have been used to densifythe composite.

12.3.1 Mixing the CNTs with the ceramic powders

A simple ultrasonic agitation in alcohol of a mixture of long CCVD–MWCNTswith nano-SiC, nano-Si3N4 or SiO2 powders seems to be insufficient todisperse the CNT aggregates, which appear on SEM images of the densifiedcomposites.9–11 The use of an ultrasonic probe, more powerful than a bath,to disperse arc-MWCNTs, shorter than CCVD ones, in nano-Al2O3 powdersis more efficient.12 To mix SWCNTs, most of which are included in ropes,with nanometric alumina, Zhan et al13,14 used, after agitation of an ethanolsuspension in an ultrasonic bath, ball-milling for 24 h with zirconia media. SEMobservations reveal only little damage to the CNTs. Wang et al.15 conductedTEM observations on powders obtained by the same preparation mode andreported good dispersion. Note that, in these works, the carbon filaments areropes of SWCNTs, at least 10 nm in diameter, and not individual SWCNTs.

A few authors have worked to improve the dispersion of CCVD-MWCNTsin oxide powders by using organic additives to modify the CNT and/or oxidesurfaces.16–18 Sun et al.16 treated the CNTs in NH3 at 600∞C for 3 h, resultingin a positive surface change, and added a small quantity of polyethyleneamine(PEI) which promotes the dispersion of the CNTs upon sonication. An aqueoussuspension of nanometric alumina with polyacrylic acid is added, givingcoated CNTs, which are themselves added to another concentrated aluminasuspension. After drying and grinding, the composite contains 0.1 wt% ofCNTs. TEM observations of the dense composites reveal a good dispersionof CNTs in the alumina matrix.16 In further works, Sun and Gao17 showedthat the control of the surface nature (basic or acidic) of MWCNTs and theuse of appropriate dispersants allow their efficient coating by surface-treatedalumina or titania powders via a heterocoagulation process. These workshave proved that the dispersion of MWCNTs in oxide powders can besignificantly improved by an appropriate functionalization and the additionof an appropriate surface agent.


12.3.2 In situ synthesis of the ceramic in CNT samples

Some authors have prepared composite powders by the synthesis of theceramic in situ around the CNT. Using short arc-made MWCNTs as templates,Hwang and Hwang19 prepared silica glass rods and then mixed them with asilica powder. More precisely, the templates were surfactant–MWCNT co-micelles, prepared by using C16TMAB (cetyltrimethylammonium bromide),on which silicates were polymerized from sodium silicate by treatment at110∞C during 48 h in an autoclave. Ning et al.18 used a sol-gel process toprepare a composite made up of very long CCVD-made MWCNTs in a silicamatrix. Several kinds of surfactants were tested – cationic (C16TMAB), anionic(polyacrylic acid) and nonionic (C16EO) with ultrasonic agitation to dispersethe CNTs. TEOS (Si(C2H5O)4) was used as raw material of SiO2. The threekinds of surfactants were efficient at dispersing the CNTs in water. The SEMstudy of the dense samples gave evidence that CNTs are more homogeneouslydispersed and have less agglomeration in composites prepared with C16TMABthan in composites prepared without any surfactant.

Huang and Gao.20 prepared MWCNT–BaTiO3 composite materials inthree steps. Firstly, rutile TiO2 particles were immobilized on the walls ofCCVD-made MWCNTs. After treatment of CNTs in nitric acid reflux (140∞C,2 h) to functionalize the walls, this was performed by, the addition of TiCl4to the suspension of functionalized CNTs at 90∞C followed by a reactiontime of 6 h. Secondly, barium acetate and NaOH were added and the reactionwas conducted under hydrothermal conditions (160∞C, 8 h). Finally, theobtained product was mixed and wet ball-milled with a BaTiO3 powder. Thedispersion of CNTs was homogeneous, owing to the better compatibilitywith the BaTiO3 powder of the MWCNTs covered by BaTiO3 than the as-prepared MWCNTs.

Recently, Liu and Gao21 prepared MWCNT–NiFe2O4 composite materialsby in situ chemical precipitation of metal hydroxides followed by hydrothermalprocessing. A long duration oxidation (10∞C, 8 h) of the CNTs by a mixtureof concentrated sulfuric acid and nitric acid was conducted to get theirfunctionalization. Then, nickel and iron nitrates were added to an ethanolsuspension of these CNTs, and after addition of sodium hydroxide (up to pH= 8.5) and a 2 h stirring, the product was treated in an autoclave at 110∞C fora few hours. In comparison with composites prepared by the same route, butusing non-oxidized CNTs, the dispersion was more homogeneous, as evidencedby SEM and TEM observations and also by an increase of the electricalconductivity of the corresponding dense materials.

Jiang and Gao22 reported the preparation of MWCNT–TiN and MWCNT–Fe2N composites using CCVD CNTs oxidized in nitric acid (140∞C, 24 h).A mixed solution of Ti(OC4H9)4 and ethanol was added to an aqueoussuspension of CNTs. The dried precursor was calcinated in N2 (450∞C, 2 h)


and then treated in flowing NH3 (800∞C, 5 h) to lead to the CNT–TiNproduct. The CNT–Fe2N product was prepared similarly using a Fe(III)–urea complex as the precursor. Jiang and Gao23 also obtained a MWCNT–magnetite composite by an in situ solvothermal synthesis from a Fe(III)–urea complex. Besides these operations to produce composite powders withthe aim of preparing dense materials, further work was devoted to coatingCNTs with in situ synthesized oxide particles, for instance via a sol-gelprocess24,25 or the precipitation of a hydroxide.26 Thus many processes areefficient at synthesizing oxide or nitride particles in situ on functionalizedMWCNTs to obtain composite powders in which the dispersion of MWCNTshas much better homogeneity than in powders obtained by mixing of MWCNTsand ceramic particles.

12.3.3 In situ synthesis of the CNTs in theceramic powder

The third way to prepare CNT–ceramic composite powders is via the synthesisof CNT by a CCVD process, in situ in the ceramic powder. A ceramicpowder which contains catalytic metal particles at a nanometric size, appropriateto the formation of CNTs, is treated at a high temperature (600–1100∞C), inan atmosphere containing a hydrocarbon or CO. In the method reported in1997 by the present authors,27 iron nanoparticles are generated in the reactoritself, at a high temperature (>800∞C), by the selective reduction in H2/CH4

(18% CH4) of an a-Al2O3 based oxide solid solution:

Al2 – 2xFe2xO3 + 3xH2 Æ (1 – x)Al2O3 + 2xFe∞ + 3xH2O

with x £ 0.10

Many clusters of Fe atoms are formed both inside and at the surface ofeach alumina grain and progressively grow. When the Fe nanoparticles locatedat the surface of the alumina grains reach a size around 0.7–2 nm, theyimmediately catalyze the decomposition of CH4 and the nucleation and growthof CNTs of very small diameter:

CH4 Æ C(Fe∞) Æ CNT(Fe/Fe3C)/Al2O3

The obtained CNT–Fe–Al2O3 powder is composed of clean isolated CNTsand small-diameter bundles of CNTs, surrounding all the oxide grains as aweb. SEM and HRTEM studies have shown that most CNTs are SWCNTs orDWCNTs (80%), with only a small proportion of three- to six-walled CNTs,and have external diameters between 0.7 and 5 nm. The method has beenextended to MgAl2O4 (Fig. 12.3) and MgO based powders, usingMg(1–x)MxAl2O4 or Mg(1–x)MxO (M = Co, Fe or Ni) as starting materials.28,29

The products which were used to prepare dense composites were CNT–Fe–


Al2O3, CNT–Co–MgAl2O4 and CNT–Co–MgO powders. In CNT–Co–MgAl2O4, most CNTs were SWCNTs. With such MgAl2O4 based powdersmodified by the additional use of a molybdenum compound, and of catalyticmaterials in the form of ceramic foams as opposed to powders, it was possibleto synthesize in situ up to 25 vol% of SWCNTs.30,31 With submicronicstarting oxide powders, the homogeneity of the distribution of CNTs insidethe matrix is very high (Fig. 12.3), probably much better than that obtainedwith other methods. However, this method presents the disadvantage of beinglimited to a few matrices, i.e. those allowing the substitution of a catalyticallyactive species in their cationic sub-lattice. It must also be pointed out that theobtained composite powders contain both a dispersion of CNTs and a dispersionof metal (and/or carbide) nanoparticles.

Another way to prepare catalytic metal nanoparticles in a ceramic powderis to impregnate the powder by a solution of a precursor salt. Weidenkaff etal.32 impregnated a substituted LnCoO3 (Ln = Er, La) powder with a citrate,

200 nm

200 nm

200 nm

200 nm

(a) (b)

(c) (d)

12.3 High resolution SEM images of SWCNT-Co-MgAl2O4 compositepowders prepared by in situ synthesis of SWCNTs within the oxidepowders, showing the high homogeneity of the distribution of CNTand the control of the CNT content. CNT contents: (a) 2.5 vol%, (b)15.0 vol%, (c) 18.3 vol%, (d) 24.5 vol%. Reprint from Acta Materilia,vol, 52, No. 4, Rul S., Lefevre-Schlick F., Capria E., Laurent Ch. andPeigney A., Percolation of single-walled carbon nanotubes in ceramicmatrix nanocomposites, pages 1061–1067, Copyright (2004) withpermission from Elsevier.


precursor or the same phase, but containing an slight excess of cobalt. MWCNTswere formed in situ by treatment of this material at 700∞C in acetylene. Anet al.33 mixed alumina, a small quantity of MgO and some iron nitrate byplanetary ball-milling in ethanol. After calcination, this material was treatedin a C2H2/H2/N2 gas flow (750∞C, 2 h) to form MWCNTs, well distributedin the alumina powder. These methods, based on the use of catalytic particleswhich are produced from the impregnation of the matrix by a precursor, offera large choice of ceramic matrices, but generally lead to MWCNTs becausecatalytic particles become too large, at a high temperature, to catalyze theformation of SWCNTs. A similar method was also used to prepare CNT–MgO composite films.34 Iron and magnesium nitrates and molybdenumacetylacetone were added to a block polymer in ethanol suspension anddeposited on a silica substrate. Treatment in CH4 (950∞C, 15 min) led to theformation of bundles of CNTs, which were identified as SWCNTs, at thefilm surface. Another interesting process has been reported by Kamalakaranet al.35 and consists of the spray pyrolysis of a slurry of g-Al2O3 and ferrocenein xylene, which was sprayed at 1000∞C using argon as the carrier gas. Theyobtain large flakes composed of a large quantity of MWCNTs intricatelymatted in a glassy alumina matrix. This continuous process could be promising,if it can be optimized to control the quantity and the quality of CNTs and toobtain a crystallized matrix.

12.3.4 Densification of CNT-ceramic composites

Hot-pressing (HP) is the most common method which has been used todensify CNT–oxide11,12,18,20,31,33,36–38 or CNT–SiC composites.9 The presentauthors showed that, by HP at temperatures between 1450 and 1530∞C,under 43 MPa and in a primary vacuum, CNT–Fe–Al2O3 nanocompositesreach a lower densification than the corresponding carbon-free Fe–Al2O3

nanocomposites, and that both the matrix grain growth and densification areinhibited when the CNT content is increased.36–38 Moreover, some CNTs aredamaged by HP at 1500∞C, producing disordered graphene layers gatheredat matrix grain junctions.38 Similar results are obtained on CNT–Fe/Co–MgAl2O4 composites, with only a densification of 90.6% for the compositescontaining 4.9 wt% of SWCNTs versus 98.2% for the corresponding Fe/Co–MAl2O4 material.38 The study of densification by HP (1300∞C, 43 MPa,secondary vacuum) of 15 different SWCNT–Co/Mo–MgAl2O4 compositescontaining between 1.2 and 16.7 vol% CNT has shown that CNTs do influencethe rearrangement step.39 For low quantities of CNT (up to 8 vol%), CNTsare favorable to this process, but for higher contents they are detrimentalbecause they form too rigid a web structure.39 In the second step, where themain active process of densification is a plastic flow controlled by a Nabarro–Herring diffusion creep,40 CNTs are as detrimental to the process when their


content increases, from 5.0 to 16.7 vol%,39 leading to a decrease in densificationfrom 95% to 73%. For CNT–Co–MgO composites (2.8 vol% of SWCNTs orDWCNTs), it is necessary to increase the HP temperature up to 1600∞C toreach a densification of 93%, which results in the destruction of most CNTs.38

The quantity of undamaged CNTs retained in the dense composite is moredependent on the treatment temperature than on the nature of the oxidematrix.38 In order to align the CNTs, the present authors have performedhigh-temperature extrusion of CNT-ceramic composites, at 1500∞C with Al2O3

and MgAl2O4 matrices and at 1730∞C for the MgO matrix.41 In the first twocomposites, CNTs were globally aligned with no more damage than in HPcomposites, showing their high resistance to the shear stress developed duringthe extrusion. But CNTs were destroyed in the MgO–matrix composite owingto the too high temperature.

An et al.33 performed HP of MWCNT–Al2O3 composites (CNTs obtainedby CCVD) at 1800∞C, in argon at a pressure of 40 MPa. The negative influenceof CNTs on the densification was reported. SEM images of fractures showedthat the grain size of Al2O3 tends to decrease when the CNT content isincreased (from 2.7 to 12.5 wt%), and that MWCNTs are located at grainboundaries. Thus, MWCNTs appear to be much more resistant at temperaturesbetween 1500 and 1800∞C than SWCNTs or DWCNTs. Siegel et al.12 reportedthe near-total densification by HP at 1300∞C in argon, under an appliedpressure of 60 MPa, of a mixture of nanophase alumina powders with 10vol% of short MWCNTs obtained by the arc method. Both the high reactivityof the matrix and the particular characteristics of the arc-made MWCNTsexplain this good result. Ning et al.11 prepared MWCNT–SiO2 materials(CNTs obtained by CCVD) by HP at 1300∞C, in N2, and with an appliedpressure of 25 MPa. The sintering mechanism was the viscous flow ofamorphous silica. These authors reported the presence of agglomerates ofCNTs and a decrease of the densification from 98% to 74% when the CNTcontent was increased from 5 to 30 vol%. In a further work,18 they showedthat enhancement of homogeneity, by using surfactant addition to preparethe starting powders, is favorable to the densification.

Huang and Gao20 studied HP, in N2 and at 25 MPa, of MWCNT–BaTiO3

composites (CNTs obtained by CCVD). They reported 1200∞C as the optimaltemperature (for 0.1 vol% CNT) and, for this HP temperature, a decrease ofthe densification from 99% to 86.5% when the CNT content increases from0.1 to 3.0 wt%, which is correlated with a much smaller matrix grain size.Ma et al.9 studied the densification by HP at 2000∞C (25 MPa, Ar atmosphere)of a mixture of 10 wt% MWCNTs, obtained by CCVD, with nanometric SiC(80 nm). The densification obtained was only 64.7%. Thus, the addition of1 wt% B4C and the increase of the temperature up to 2200∞C were necessaryto reach a densification of 98.1%. These authors reported some images showingthat at least some of the MWCNTs had not been destroyed by this very high


temperature treatment. The only work of hot-isostatic pressing (HIP) ofCNT–ceramic composites so far was reported by Balazsi et al.,10 and used amixture of 1 wt% CCVD-MWCNTs with Si3N4, treated at 1700∞C. Theyshowed that MWCNTs remained after HIP at 2 MPa for 1 h, but that theyhad completely disappeared after treatment at 20 MPa for 3h.

An efficient method to achieve the total densification of CNT–ceramiccomposites without damaging the CNTs could be the spark plasma sintering(SPS) technique. Zhan et al.13 studied the SPS of a mixture of ropes of 5.7or 10 vol% SWCNTs (obtain by CCVD) with nanocrystalline alumina(40 nm), containing both the a and g forms. Relative densities of 100% bySPS at 1150∞C, for only 3 min and under an applied pressure of 63 MPa,were obtained for pure alumina and for the two CNT–Al2O3 composites.Owing to the low temperature reached and the very short time of treatment(heating within a few minutes), CNTs were not damaged. The SPS techniqueis well known to combine the effects of rapid heating, pressure and powdersurface cleaning and that could explain its efficiency for CNT-ceramic materials.An appropriate microstructure of the powder, i.e. highly reactive aluminaand SWCNTs in ropes rather than individual, probably also favors a betterdensification. Wang et al.15 sintered the same materials using SPS, but at1450–1550∞C under an applied pressure of 40 MPa, resulting in no moredamage of CNTs than in the former work13 but in lower densifications (95.1%for 10 vol% CNTs). Sun et al.16 reported the full densification by SPS at1300∞C (5 min) of a MWCNT–Al2O3 mixture prepared by a colloidal process,but the result is less significant because of the very low CNT content (0.1wt%).

The difficulty of achieving the densification of CNT–ceramic compositesis a critical problem, particularly when good mechanical properties are essential.Generally, the CNTs inhibit the matrix grain growth and the densificationprocesses. Hot-pressing is often not efficient to totally densify the materials,particularly when the CNT content is higher than a few vol%. Higherhomogeneity of the CNT distribution, such as that obtained by the in situformation of CNTs, seems to make the problem worse. Increasing thetemperature treatment and/or the applied pressure can contribute to increasingthe densification, but leads to damage to the SWCNTs or DWCNTs. However,large-diameter MWCNTs are generally much more resistant. In comparisonwith HP, the SPS method involves shorter treatments and lower temperaturesto obtain a complete densification, when the microstructure of the greenmaterial is optimized. Thus SPS is attractive because it avoids or at leastlimits the damage to the CNTs due to high temperatures.


12.4 Properties of CNT-ceramic composites


Most research on CNT-ceramic composite materials has for its main purposebeen directed at the reinforcement of the ceramic, i.e. obtaining a significantincrease in the fracture toughness of these brittle materials. Most authorshave also measured the fracture strength because a tougher material withlower fracture strength is not desirable. Microscopy studies were also conductedon damaged or fractured materials to investigate the behavior of the CNTsincluded in the ceramic matrix, with the aim of identifying possiblereinforcement mechanisms. These works were conducted mainly with oxidematrices such as Al2O3,

12,13,15,16,36,38,42,43 MgAl2O4,38 and glassy or crystallized

SiO2,11,18,19 and a only few were related to SiC9 or Si3N4.10 Two groups havereported the tribological properties of CNT–alumina composites33 and ofcarbon/carbon composites covered by a film of CNTs.44

The first results on the mechanical properties of alumina–matrix compositescontaining CNTs were reported in 1998 by the present authors.36 SeveralCNT–Fe–Al2O3 composite powders with different CNT contents (0.7–6 wt%of carbon, mainly SWCNTs or DWCNTs) were prepared by reducingAl2(1–x)FexO3 solid solutions of different compositions (x = 0.02–0.20) inH2–CH4, at 900 or 1000∞C. The hot-pressed composites had fracture strengths(measured by three-point bending) lower than the corresponding Fe–Al2O3

nanocomposites45 and fracture toughness (measured by the single-etchednotched beam (SENB) technique) similar to or slightly lower than that ofAl2O3 (3–5 MPa.m1/2 and 4.4 MPa.m1/2 respectively). The matrix grainswere micrometric and the CNTs were located either at grain boundaries or atgrain junctions, or within the grains (Fig. 12.4(c)). The pullout of someCNTs was evidenced, showing a possible reinforcement mechanism (Fig.12.4(c)). The absence of real reinforcement was attributed to the uncompleteddensification of the materials, a too low CNT volume fraction, and the presence,in some composites at least, of undesirable forms of carbon (carbon nanofibers,nano-ribbons). In a further work,37 an attrition-milled starting oxide wasused to enhance the homogeneity of the distribution of the CNTs within thematrix (100–500 nm grains – Figs 12.4(a), (b)), but both the fracture strengthand fracture toughness were lower than that of the previous materials. Then,the quality of CNT was improved and the preparation method was extendedto MgAl2O4 and MgO matrices, without obtaining any improvement of thesemechanical properties, in comparison with the pure matrix.38 As reported inthe previous section, whatever the nature of the oxide matrix, CNTs inducea low densification of the HP materials, and this effect increases with theCNT content, resulting in relatively poor mechanical properties.

Siegel et al.12 reported, in 2001, an improvement of the fracture toughnessfrom 3.4 MPa.m1/2 for hot-pressed Al2O3 to 4.2 MPa.m1/2 for a hot-pressed


MWCNT–Al2O3 composite (10 vol% of short MWCNTs prepared by the arcmethod). Note that the fracture toughness was calculated by using the Evansand Charles equation, from the lengths of cracks emanating from Vickersindentations (5 kg load). Sun et al.16 reported a fracture toughness (determinedby the same method) of 4.9 MPa.m1/2 for a MWCNT–Al2O3 composite fullydensified by the SPS method and containing only 0.1 wt% of CNTs (longMWCNTs prepared by CCVD), in comparison to 3.7 MPa.m1/2 for Al2O3.These authors underlined the tight bonding between the CNTs and the matrixand inferred a bridging effect on cracks and some pullout as possiblereinforcement mechanisms.

The results which have attracted most attention, however, were those

(a) 200 nm

(b)(c)

100 nm 100 nm

12.4 SEM images of fracture surface of CNT-Fe-Al2O3 compositesdensified by hot-pressing of composites powders, which have beenprepared by in situ synthesis of CNTs (mainly SWCNTs and DWCNTs)within the oxide powder. (a) and (b) specimen prepared by usingsubmicronic oxide powders, (c) specimen prepared by usingmicronic oxide powders. Reprint from Ceramic International, vol. 26,No. 6, Peigney A., Laurent Ch, Flahaut E. and Roussel A., Carbonnanotubes in novel ceramic matrix nanocomposites, pages 1061–1067, Copyright (2000) with permission from Elsevier.


reported by Zhan et al.13 on SWCNT–Al2O3 composites, with an increase ofthe fracture toughness (determined from Vickers indentations using the Antisequation) from 3.3 MPa.m1/2 for Al2O3 to 7.9 MPa.m1/2 for 5.7 vol% CNTs,and up to 9.7 MPa.m1/2 for 9.7 vol% CNTs. The hardness simultaneouslydecreased from 20.3 to 20.0 GPa, and then to 16.1 GPa. Remember that thesecomposites were fully densified by SPS at 1150∞C, and that the startingmaterials were ropes of SWCNTs (obtained by CCVD) mixed by ball-millingwith a nanocrystalline Al2O3 powder (40 nm), containing both the a and gforms. The ropes of SWCNTs are located at intergranular positions andseemed to be undamaged (Fig. 12.5).

Very recently, Wang et al.15 reported the results of measurements, by thesingle-edged V-notch beam (SEVNB) method, of the fracture toughness ofSWCNT–Al2O3 composites similar to those of the previous reference:13 usingthe same starting material, with 10 vol% SWCNTs, but sintering by SPS at1550∞C instead of 1150∞C, giving a densification of 95.1%. They reported afracture toughness of only 3.32 MPa.m1/2, similar to that of Al2O3 or graphite–Al2O3 composites (3.22 and 3.51 MPa.m1/2 respectively). The ropes of SWCNTswere not damaged by the short SPS treatment at 1550∞C but the matrixgrains seemed slightly larger and the densification lower than that obtainedby SPS treatment at 1150∞C.13 By Hertzian indentations using a tungstencarbide ball, they showed that no Hertzian cracks occur on such composites,in the conditions in which they are widely formed on alumina. So, CNT–

1000 nm

12.5 SEM image of the fracture surface of 5.7 vol% SWCNT-Fe-Al2O3composite densified by spark plasma sintering (SPS) of a mixture ofnanometric alumina and ropes of SWCNTs. Reprint from NatureMaterials, No. 2, 2002, pp. 38–42, Zhan G.-D., Kuntz J.D., Wan J. andMukherjee A.K., ‘Single-wall carbon nanotubes as attractivetoughening agents in alumina-based nanocomposites’, with thepermission of Nature Materials and of the authors (http://www.nature.com/nmat/index.html).


alumina composites are highly contact-damage resistant, similarly to graphite–alumina composites, and could be used in all applications where contactloading is important. Moreover, these results showed that the fracture toughnessof such composites can be severely overestimated when measured by thestandard indentation method.46 Indeed, so far neither the SENB nor theSEVNB results have evidenced that CNT can significantly reinforce aluminaceramics.

Hwang and Hwang19 studied the hardness of glassy SiO2 in which eitherSiO2 rods or CNT–SiO2 rods (synthesized using MWCNTs as templates) aredispersed. The CNT–SiO2 rod–SiO2 materials had a significantly higherhardness than the SiO2 rod–SiO2 composites. However, the real influence ofCNT on the hardness was not proved, because the presence of CNT insideeach rod in the dense composite is questionable. Probably, the higher hardnessis a consequence of a better morphology of the rods prepared by using theCNT templates. Thus, using the words ‘Carbon nanotube reinforced ceramics’as the title of the paper19 is debatable. Ning et al.11 reported an increase ofbending strength from 50 to 85 MPa (three-point bending) in 5 vol% MWCNT–SiO2 composites densified at about 98%, in spite of the fact that the homogeneityof the distribution of CNT seemed not ideal (Figs 12.6(a), (b)). They alsoreported an increase of the fracture toughness from 1 to 2 MPa.m1/2, butthese values were determined by the indentation method, and thus the secondvalue could be overestimated. Moreover, CNTs provide very suitable conditionsfor the nucleation and crystallization of cristobalite,11 and that could contributeto the evolution of the mechanical properties. In further work, Ning et al.18

enhanced the homogeneity of the CNT distribution (see the previous section)and obtained full densification of the 5 vol% CNT-containing composite;they reported still higher bending strength (97 MPa) and fracture toughness(2.46 MPa.m1/2). From TEM images, they inferred that a good interfacebonding exists between the MWCNTs and glassy SiO2. These results onglassy materials are to be considered, but the fracture toughness has to beconfirmed by measurements using a SENB-type method.

Ma et al.9 reported 10% increases, both in fracture strength (three-pointbending) and in fracture toughness (SENB method), in MWCNT–SiCcomposites hot-pressed at 2000 or 2200∞C. The homogeneity of the distributionof CNTs was low and the increases were too small to be significant. Balazsiet al.10 compared the elastic modulus and the bending strength of severaldifferent materials prepared by HIP: Si3N4 alone and Si3N4 with additions ofcarbon black (23 wt%), graphite (23 wt%), carbon fibers (1 wt%) or MWCNTs(1 wt%) which all decreased the apparent density of the material. The highermodulus and strength values were obtained for Si3N4 alone and, although thevalues decreased less with the addition of MWCNTs than with addition ofother additives, these results did not justify inclusion of the words ‘carbonnanotube reinforced silicon nitride composites’10 in the title of the paper.


Very recently, Xia et al.43 reported microstructural investigations onMWCNTs which had been formed within the regular and well-aligned poresof an alumina membrane. The material was too thin (20–90 mm) to permitmechanical measurements, but different possible reinforcement mechanismsinduced by the CNTs were evidenced on stressed and damaged materials,

(a)

(b)

12.6 SEM images at low (a) and high magnification (b) of surfacefractures of 5 vol% MWCNT-SiO2 composites densified at about 98%by hot-pressing, showing a rather low degree of homogeneity of theCNTs distribution. Reprint from Materials Science & Engineering, A:Structural Materials: Properties, Microstructure and Processing A,Vol. 357, No. 1–2, Ning J., Zhang J., Pan Y. and Guo J., Fabricationand mechanical properties of SiO2 matrix composites reinforced bycarbon nanotube, pages 392–396, Copyright (2003) with permissionfrom Elsevier.


such as crack deflection, crack bridging and CNT pullout. Moreover, a newmechanism of CNT collapse in shear bands occurs (Fig. 12.7), rather thancrack formation, suggesting that these materials can have a multiaxial damagetolerance.43 These results showed that the key problem to solve in order toobtain real reinforcement of a ceramic is probably the preparation ofmacroscopic samples in which the CNTs would have been well organizedwithin the matrix.

An et al.33 hot-pressed MWCNT–Al2O3 composites at 1800∞C and notedboth a decrease of the matrix grain size when the CNT content increases anda poor cohesion between the CNTs and alumina. For CNT contents up to 4wt%, the microhardness was enhanced and the wear loss decreased, maybeowing to a lower matrix grain size. At higher CNT contents, the evolution ofthese parameters was reversed. Thus, the influence of the CNTs was notclearly established. Lim et al.44 applied on carbon/carbon composites differentcoatings made of a carbon-based material containing MWCNTs (synthesizedby a CCVD method). They showed that the wear loss decreases regularly (upto 100% for 20 wt% CNT) but also that the friction coefficient increasesslightly when the CNT content in the coating increased, showing a significantinfluence of the CNTs on tribological properties.

12.4.2 Electrical properties

Only a few authors have reported results concerning the electrical propertiesof CNT-ceramic composites and most have been concerned with the influence

12.7 SEM images of the deformation around an indentation in a90 mm thick MWCNT-Al2O3 composite prepared by in situ formationof CNTs within the regular and very well aligned pores of an aluminamembrane. (a) array of the ‘shear bands’ formed, (b) close up viewof lateral buckling or collapse of the CNTs in one ‘shear band’.Reprint from Acta Materialia, Vol. 52, Xia Z., Ricster L., Curtin W.A.,LiH., Sheldon B.W., Liang J., Chang B. and Xu J.M., Directobservation of toughening mechanisms in carbon nanotube ceramicmatrix composites, pages 931–944, Copyright (2004) with permissionfrom Elsevier.

(a) (b)


of CNT on DC electrical conductivity measured at room temperature. In2000, Flahaut et al.38 reported that CNTs confer an electrical conductivity toceramic–matrix composites, whereas the corresponding ceramics and metal–oxide nanocomposites are insulators. Owing to the percolation of the CNTs(mainly SWCNTs or DWCNTs), the DC electrical conductivity jumped fromvalues lower than 10–10 S.cm–1 for Fe–Al2O3 of Fe/Co–MgAl2O4 compositesto 0.4–4.0 S.cm–1 for CNT–Fe–Al2O3 composites, or to 1.5–1.8 S.cm–1 forCNT–Fe/Co–MgAl2O4 composites.38 These values are fairly well correlatedto the relative quantity of CNTs. For CNT–Co–MgO composites, the valuewas lower (0.2 S.cm–1) because an important proportion of CNTs were damagedor destroyed during the hot-pressing at 1600∞C.

Rul et al.31 prepared, by in situ synthesis, 22 different SWCNT–Co/Mo–MgAl2O4 composites with a wide range of CNT content (between 0.11 and24.5 vol%). Their DC electrical conductivity jumped from 10–10 S.cm–1 to0.0040 S.cm–1 between 0.23 and 1.16 vol% and reached 8.5 S.cm–1 for thehigher CNT content (Fig. 12.8(b)). It was shown that the electrical conductivitywas well fitted by the scaling law of the percolation theory, s = s0(p – pc)

t

with a percolation at a low threshold, pc = 0.64 vol%, and an exponent t =1.73 close to the theoretical value (t = 1.94) characteristic of a three-dimensionalnetwork (Fig. 12.8(a)). The low percolation threshold is a consequence ofthe very large aspect ratio (>10 000) of the SWCNT. These results showedthat the electrical conductivity of such composites can be tailored in a widescale (10–2–10 S.cm–1) through the CNT content. Peigney et al.41 showedthat the alignment of CNTs within such matrices, obtained by hot extrusion,can lead to an important anisotropy of the electrical conductivity: 20 S.cm–1

in the extrusion direction versus 0.60 S.cm–1 in the transverse direction.Zhan et al.14 also measured the DC electrical conductivity of Al2O3-basedcomposites containing ropes of SWCNTs. The DC electrical conductivityincreased up to 33.45 S.cm–1 upon increasing the CNT content, a valuesignificantly higher than that reported by Rul et al.31 with unorganized CNTs.Less damage to CNTs during the SPS sintering (1150∞C for only a fewminutes) than during hot-pressing (1300∞C, heating rate 10∞C.min–1) couldexplain this result.

Some authors have reported results on the influence of CNTs on theelectrical properties of non-insulating mixed oxides. Huang and Gao20 densifiedMWCNT–BaTiO3 composites by hot-pressing (1200∞C, 35 MPa) of BaTiO3

powders containing MWCNTs covered by in situ synthesized BaTiO3 particles.The electrical conductivity decreased when the CNT content increased (from6.9 S.cm–1 for BaTiO3 to 3.6 S.cm–1 for the composite containing 3 wt%CNT), and the n-type semi-conductivity of BaTiO3 was converted to p-type,even for a very low CNT content (0.1 wt%). This phenomenon has beenattributed to a Schottky barrier constructed at the CNT–matrix contact, whichcould be promising for the fabrication of new-style ferroelectric and


12.8 DC electrical conductivity of SWCNT-Co-MgAl2O4 compositesversus the CNT contents. (a) up to 10 vol% CNTs, the values are wellfitted by the scaling law of the percolation theory s = so(p – pc)

t witha percolation at a low threshold, pc = 0.64 vol%, and an exponentt = 1.73. (b) the value jumps from 10–10 S.cm–1 to 0.0040 S.cm–1

Reprint from Acta Materialia, vol. 52, No. 4, Rul S., Lefèvre-Schlick F.,Capria F., Laurent Ch. and Peigney A., Percolation of single-walledcarbon nanotubes in ceramic matrix nanocomposites, pages 1061–1067, Copyright (2004) with permission from Elsevier.

0 5 10 15 20 25CNT (vol %)

(b)

10

1

0.1

0.01

0.001

s (S

/cm

)

0 2 4 6 8 10CNT (vol %)

(a)

10+2

1

10–2

10–4

10–6

10–8

10–10

s (S

/cm

)

0.64

s = k (p – pc)t with pc = 0.64 ± 0.02

t = 1.73 ± 0.02

–0.5 0 0.5 1Log10 (p – 0.64)

0.5

0

–0.5

–1

–1.5

–2

–2.5

Log

10 s

(s

in S

/cm

)

Log10 s = 1.73 log10 (p – 0.64)– 1.758 R2 = 0.9921


thermoelectric devices.20 Liu and Gao21 prepared MWCNT–NiFeO4 compositesand studied the effect of a surface oxidation treatment of MWCNTs. Theincrease of electrical conductivity of the composites was larger by using pre-oxidized CNTs (from 6.8 ¥ 10–4 to 82.2 S.cm–1 versus to 9.1 S.cm–1, for 10wt% CNTs) and this was attributed to a better homogeneity of the distributionof CNTs within the matrix and also to a stronger adhesion between the CNTsand the matrix.21

Jiang and Gao23 showed that, in composites prepared by in situ synthesisof magnetite on MWCNTs produced by CCVD, the electrical conductivityincreased from 1.9 S.cm–1 for the material without CNTs to 2.5 S.cm–1 forthe material containing 32.95 wt% of CNTs. This increase seems moderate,taking into account both the large content and the electrical conductivity ofCNTs, showing that the influence of interfaces is probably determinant inthe conduction mechanisms of the material. Huang et al.47 embedded MWCNTs(arc-discharge, 3 wt%) into Bi2Sr2CaCu2O8+d (a Bi-2212 superconductor) bya partial melting processing and measured an insulator–superconductortransition slightly lower in the composite than in the Bi-2212 material (87 Kversus 95 K) and also an increase in current densities.

Owing to the percolation of CNTs in insulating ceramics, materials witha DC electrical conductivity directly tailored in a wide range (0.01 to 10–100S.cm–1) by the quantity of CNTs can be prepared. The high aspect ratio ofCNTs allows percolation thresholds lower than 1 vol%, i.e. CNT contents forwhich full densification of the composites is easily obtained. However, toreach the higher conductivity values (> 10 S.cm–1), higher CNT contents(>10 vol%) are necessary, which requires the use of the SPS method toachieve the densification of the materials. One of the critical points is thehomogeneity of the distribution of CNTs, and the in situ synthesis methodsare promising in that respect. The other critical point is the characteristic ofthe interfaces between CNTs and the matrix, which can be greatly influencedby oxidative or more complex fuctionalization treatments. In a semiconductingoxide matrix, CNTs can either increase or decrease the electrical conductivity,as a function of the conducting mechanisms involved both at the interfacesand in the matrix. As shown in ferroelectric or superconductor ceramic matrices,the addition of CNTs could notably influence the transport properties ofmany ceramics used in electronics, and possibly lead to materials with novelor improved functional properties. The organization of CNTs within thematrix makes possible the preparation of materials with anisotropic functionalproperties.

12.4.3 Thermal properties

The very high thermal conductivity of CNTs (2000–6000 W.m–1K–1) lets oneenvision that they could be used to manage the thermal properties of ceramics.


Yowell48 reported the preparation by tape-casting of films of partially stabilizedzirconia containing 1 wt% SWCNTs. However, the SWCNTs did not survivethe thermal treatment, and thus the slight increase of the measured thermalconductivity could have been due to the presence of residual metal catalystparticles.48 In contrast, the dispersion of 1 wt% vapor-grown carbon fibers(VGCF), which withstand the sintering, provoked a decrease of about 30%of the thermal conductivity. Seeger et al.49 incorporated MWCNTs into SiO2

by partial matrix melting caused by a laser. They reported that the presenceof CNTs was crucial for the heat absorption and melting of the matrix,showing that the thermal transport of SiO2 was probably greatly enhanced byCNTs. Thus, further works are required to investigate more precisely theinfluence of CNTs on the thermal properties of materials and the correspondingheat transport mechanisms. If these studies confirm that CNTs can significantlyincrease the thermal conductivity of ceramic material, the crucial point willbe to determine for which CNT content this could be achieved.

12.5 Conclusions and future trends

Because the first reports on CNT-ceramic composites date only from 1998,and because only a few teams have worked so far on these novel materials,it could be argued that we are at the infancy of the development of a newclass of composite materials. Researches on these materials depend firstly ona better knowledge of the CNTs by their users. Depending on theirmicrostructural characteristics (SWCNTs, individual or in ropes, MWCNTs,diameter, length, number of walls), but also on the synthesis methods whichhave been used, the properties of CNTs may greatly vary. Notably, thetreatments involved in the control of the surface properties and reactivity ofthe CNTs need to be optimized for a particular form of CNTs synthesized bya particular method.

To obtain good homogeneity by mixing CNTs with a ceramic powder, itis necessary to adapt the surface properties of both the CNTs and the ceramicparticles, which will be preferably nanometric, and this requires the additionof organic additives. The in situ synthesis of ceramic onto CNTs also requiresa fine and stable dispersion of CNTs in a suitable medium. The latter methodcan ensure a good adhesion between the CNTs and the ceramic, but is morecomplex and less flexible than the previous one. The in situ synthesis ofCNT within the ceramic powder leads to very homogeneous materials butrequires an in-depth knowledge of the CCVD process to avoid the formationof undesirable forms of carbon, and can be applied to only a few matrices.Thus, for each particular variety of CNT-ceramic composites, and with aview to developing a particular property, a dedicated preparation methodcould be preferred. To densify such materials, spark plasma sintering (SPS)has proved to be efficient, leading to good densifications with limited damage


to the CNTs. The possibility of increasing the fracture toughness of ceramicsthrough the addition of CNTs is still debatable. But it has been demonstratedthat for these materials, SENB-type tests are to be preferred to indentation-derived measurements which are unsuitable.

The contact damage resistance of CNT-ceramic composite could bepromising, but the use of CNTs instead of graphite will have to be justifiedby gains in the other mechanical properties. However, at the microscopicscale, several reinforcement modes have been evidenced, particularly whenthe CNTs are well aligned within the matrix. In spite of its very high aspectratio and its intrinsic mechanical properties, the interfacial surface areadeveloped by one CNT is small, due to its nanometric diameter. Thus, theeffect of only one (or a few) isolated CNT(s) is probably too low to reallyinfluence the propagation of a critical crack. But the effect of many CNTsoperating simultaneously could be useful, which requires a sufficient CNTcontent, and especially a high degree of organization of the CNTs to ensurea synergistic effect.

Owing to the high aspect ratio of CNTs which facilitates their percolation,their addition to insulating ceramics is efficient in giving electronic conductivityto the material and in tailoring the conductivity value directly by the quantityof CNTs. Moreover, anisotropic conductivity is obtained when the CNTs arealigned within the composite. In semiconducting ceramics, CNTs couldnotably influence the transport properties of many ceramics used in electronics,and possibly lead to materials with novel or improved functional properties.The thermal properties of ceramics can be modified, and the thermalconductivity possibly enhanced, via the addition of CNTs. But, for mostproperties, the organization of the CNTs within the matrix is a prerequisiteto taking full advantage of their dispersion in ceramic materials. While theviability of structural CNT-ceramic composites may be a long shot, thepromising results reviewed here give strong hope that functional materialswith tailored electrical/thermal characteristics will find their way into industrialapplications.

12.6 Sources of further information

The Nanotube Site (http://www.nanotube.msu.edu/) has general informationabout CNTs, links to sources of nanotubes and nanotube-based products andto sites relevant to nanotube research. Note that research in the carbon nanotubefield is progressing at a very fast pace (about 10 reviewed papers each day).

Another source of information is the website of the research group ofthe authors, the Nanocomposites and Carbon Nanotubes Group(http://ncn.f2g.net/).

Key books on carbon nanotubes are:


∑ Dresselhaus, M.S., Dresselhaus, G., Eklund, P.C. and editors, Science ofFullerenes and Carbon Nanotubes, Academic Press, San Diego, 1996.

∑ Harris P.J.F., Carbon Nanotubes and Related Structures – New Materialsfor the Twenty-first Century, Cambridge University Press, Cambridge,1999.

12.7 References

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(London), 1993, 363, 603–605.3. Bethune, D.S., Kiang, C.H., de Vries, M.S., Gorman, G., Savoy, R., Vazquez, J. and

Beyers, R., ‘Cobalt-catalyzed growth of carbon nanotubes with single-atomic-layerwalls’, Nature (London), 1993, 363, 605–607.

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20. Huang, Q. and Gao, L., ‘Manufacture and electrical properties of MWCNT/BaTiO3

nanocomposite ceramics’, J. Mater. Chem., 2004, 14, 2536–2541.21. Liu, Y. and Gao, L., ‘A study of the electrical properties of carbon nanotube–NiFe2O4

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22. Jiang, L. and Gao, L., ‘Carbon nanotubes–metal nitride composites: a new class ofnanocomposites with enhanced electrical properties’, J. Mater. Chem., 2005, 15(2),260–266.

23. Jiang, L. and Gao, L., ‘Carbon nanotubes–magnetite nanocomposites from solvothermalprocesses: formation, characterization, and enhanced electrical properties’, Chemistryof Materials, 2003, 15, 2848–2853.

24. Seeger, T., Kohler, T., Frauenheim, T., Grobert, N., Rühle, M., Terrones, M. andSeifert, G., ‘Nanotube composites: novel SiO2 coated carbon nanotubes’, ChemicalCommunications, 2002, 34–35.

25. Vincent, P., Brioude, A., Journet, C., Rabaste, S., Purcell, S.T., Brusq, J.L. andPlenet, J.C., ‘Inclusion of carbon nanotubes in a TiO2 sol-gel matrix’, J. Non. Cryst.Solids, 2002, 311, 130–137.

26. Hernadi, K., Ljubovic, E., Seo, J.W. and Forro, L., ‘Synthesis of MWNT-basedcomposite materials with inorganic coating’, Acta Materialia, 2003, 51, 1447–1452.

27. Peigney, A., Laurent, Ch., Dobigeon, F. and Rousset, A., ‘Carbon nanotubes grownin situ by a novel catalytic method’, J. Mater. Res., 1997, 12, 613–615.

28. Govindaraj, A., Flahaut, E., Laurent, Ch., Peigney, A., Rousset, A. and Rao, C.N. R.,‘An investigation of carbon nanotubes obtained from the decomposition of methaneover reduced Mg1–xMxAl2O4 spinel catalysts’, J. Mater. Res., 1999, 14, 2567–2576.

29. Flahaut, E., Peigney, A., Laurent, Ch. and Rousset, A., ‘Synthesis of single-walledcarbon nanotube–Co–MgO composite powders and extraction of the nanotubes’, J.Mater. Chem., 2000, 10, 249–252.

30. Rul, S., Laurent, Ch., Peigney, A. and Rousset, A., ‘Carbon nanotubes prepared insitu in a cellular ceramic by the gelcasting-foam method’, J. Eur. Ceram. Soc., 2003,23, 1233–1241.

31. Rul, S., Lefèvre-Schlick, F., Capria, E., Laurent, Ch. and Peigney, A., ‘Percolationof single-walled carbon nanotubes in ceramic matrix nanocomposites’, Acta Mater.2004, 52, 1061–1067.

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33. An, J.W., You, D.H. and Lim, D.S., ‘Tribological properties of hot-pressed alumina–CNT composites’, Wear, 2003, 255, 677–681.

34. Du, C. and Pan, N., ‘Preparation of single-walled carbon nanotube reinforced magnesiafilms’, Nanotechnology, 2004, 15, 227–231.


35. Kamalakaran, R., Lupo, F., Grobert, N., Lozano-Castello, D., Jin-Phillipp, N. Y. andRühle, M., ‘In-situ formation of carbon nanotubes in an alumina–nanotube compositeby spray pyrolysis’, Carbon, 2003, 41, 2737–2741.

36. Laurent, Ch., Peigney, A., Dumortier, O. and Rousset, A., ‘Carbon nanotubes–Fe–alumina nanocomposites. Part II: Microstructure and mechanical properties of thehot-pressed composites’, J. Eur. Ceram. Soc., 1998, 18, 2005–2013.

37. Peigney, A., Laurent, Ch., Flahaut, E. and Rousset, A., ‘Carbon nanotubes in novelceramic matrix nanocomposites’, Ceram. Int., 2000, 26, 677–683.

38. Flahaut, E., Peigney, A., Laurent, Ch., Marlière, Ch., Chastel, F. and Rousset, A.,‘Carbon nanotube–metal-oxide nanocomposites: microstructure, electrical conductivityand mechanical properties’, Acta Mater., 2000, 48, 3803–3812.

39. Peigney, A., Rul, S., Lefèvre-Schlick, F. and Laurent, Ch., ‘Densification during hot-pressing of carbon nanotube–metal–magnesium aluminate spinel composites’, to besubmitted, 2006.

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334

13.1 Introduction

Advanced ceramic materials have been successfully developed over the pastfew decades. They are widely used in a variety of applications for theirsuperior properties in comparison to conventional materials, such as metalsand polymers. A few of these properties include high strength-to-mass ratio,excellent wear resistance and exceptional corrosion resistance.

Although the development of advanced ceramic materials has progressedtremendously over the years, barriers to their wide acceptance exist. One ofthese barriers is the inherent difficulty in materials processing. For instance,after sintering, ceramic materials become hard and brittle. In the processingof ceramic materials using traditional methods, such as machining and grinding,fracture occurs at stress-concentration locations which leaves cracks on andbeneath the machined surfaces of ceramic components. These processing-induced damage areas degrade the quality of products and often lead tomalfunction and/or catastrophic failure during the service life.

Research on machinable advanced ceramics has concentrated on multiphasecomposites and materials design of ceramics. Most of these research effortsaim at improving the ceramic machinability while maintaining the material’sdistinct features. Recently, nanocomposites in which nano-sized particleswere dispersed within the matrix grains or at the grain boundaries showmuch better mechanical properties (hardness and strength) as well asmachinability and superplasticity compared with monolithic materials. Inthis chapter, some machinable microcomposites and nanocomposites withweak or soft phase (h-BN, LaPO4 and Ti3SiC2) will be discussed based onthe relationship between properties and microstructure.

13.2 Design principles of machinable ceramics

Ceramic composites may be defined as those ceramic materials that consistof two or more fundamentally different components that are able to act

13Machinable nanocomposite ceramics

R W A N G, Arizona State University, USA

Machinable nanocomposite ceramics 335

synergistically to give properties superior to those provided by either componentalone. Compound machinable ceramics are defined as those having a distinctweak interface phase or layered phase distributed throughout the bulk matrixceramic. By varying the type and distribution of the weak interface phase orlayered phase in the composite, it is possible to obtain a wide range ofmechanical properties and machinability combinations. Such ceramic materialshave a number of potential advantages for fabrication of complex-shapeengineering components. Table 13.1 gives some basic materials propertiesfor advanced ceramics and layered or soft materials. Based on thethermodynamic properties and sintering compatibility of these advancedceramics and layered materials, a variety of machinable ceramic compositescan be designed and constructed, such as Si3N4/h-BN, SiC/h-BN, AlN/h-BN, Al2O3/h-BN, Al2O3/LaPO4, Al2O3/Ti3SiC2, Ce–ZrO2/CePO4, etc.

Unreactive or weak bonding is the main design principle of compoundmachinable ceramics. The addition of a layered or weak ceramic used asboundary phase in the ceramic–matrix composite can improve the machinabilityof advanced ceramics by deflecting crack propagation. The following energy-absorbing mechanisms have been identified from SEM or TEM images ofareas in the vicinity of the indentation: diffuse microcracking, delamination,crack deflection, grain pullout and the buckling of individual grains. All ofthese might be attributed to the improved machinability of composites. Figure13.1 shows the crystal structure and morphological microstructure of h-BN,Ti3SiC2 and LaPO4, so-called weak or layered phase. You can see that thesematerials possess a layered structure or step-way fracture characteristics,and form weak bonding with the matrix ceramics. Based on this designprinciple of machinable ceramics, oxide (Al2O3/LaPO4) and nonoxide (Si3N4/h-BN) have been investigated and introduced considering the relationshipbetween the mechanical properties, microstructure and machinability. In someresearch, nanocomposites showed potential machinability and superplasticity[1], as we will also discuss in this chapter.

13.3 Al2O3–LaPO4

Alumina (Al2O3) has been recognized as one of the most promising structuralmaterials for many mechanical or thermomechanical applications because ofits excellent high-temperature strength, good oxidation and wear resistance,high hardness, and low specific weight. Because of its strong covalent bondingcharacter, extremely hard Al2O3 ceramics make conventional machining verydifficult or even impossible. The addition of a weak interface phase or layeredphase in the matrix to facilitate crack deflection and propagation duringmachining, named compound machinable ceramic, was used for improvingthe machinability of ceramic such as mica glass-ceramic. Lanthanum phosphate(LaPO4), also known as monazite, has been found to be a suitable and effective

Ceram

ic matrix com

posites336

Table 13.1 Basic material properties for advanced ceramics and layered or soft materials

Advanced ceramics Layered or soft materials

Si3N4 SiC Al2O3 Y-PSZ C h-BN Ti3SiC2 Mica LaPO4

Density (g cm–3) 3.19 3.22 3.98 6.1 2.265 2.27 4.5 2.889 5.07

Hardness (GPa) 14–18 21–25 19.3 8–12 1–2* 2* 3–5 2–4 4~7

Elastic modulus 280–320 450 400–410 180–220 2.5–10 50–80 320 — 100–220(GPa)

Bending strength 400–1000 640 550–600 650–1000 20–70 40–60 >250 170–360 90–140(MPa)

Fracture toughness, 3.4–8.2 5.7 3.8–4.5 6–8 <1 7 — ~1.0KIC (MPa.m1/2) RT RT

Thermal expansion 2–3 4.2 6.5–8.9 10.6 1.2–28.3 0.59–10.51 8.6–9.7 — 10coefficient 0–1000∞C 0–1500∞C 200–1200∞C 0–1000∞C 20–800∞C(¥ 10–6C–1)

Thermal 15–50 50 38.9 1–2 83.7–272 40–70 20–40 — 5conductivity(W m–1 K–1) 600∞C 20∞C

*Moh’s scale.


debonding material for high-temperature oxide/oxide composites [2]. Fractureenergy measurements have shown that the Al2O3/LaPO4 interface was weakenough, when compared to the Al2O3/Al2O3 interface, to satisfy the He andHutchison debonding criteria [3] for bimaterial interfaces.

Recent work by Davis and Marshall [4] has shown monazite to be asuitable weak bonding material in preparing composites with Al2O3. Thesecomposites show high toughness with an interface sufficiently weak fordebonding to occur when cracks are deflected into monazite and away fromthe Al2O3 phase. LaPO4 in the Al2O3 composites is quite stable and noreaction occurs between the two phases up to 1600∞C, provided the La:Pratio in the monazite is close to 1. Furthermore, composites of LaPO4 withsome other ceramics are machinable as evaluated from information availablein the literature. Interface properties of LaPO4 with Al2O3 and yttrium aluminum

B atomN atoma0 = 2.054 Å

C0 = 6.660 Å

Signal A = InLensDate: 28 Sep 2001

EHT = 10.00 kV WD = 8 mm

2mmMax = 20.00 KX

(a) (b)

2mmMag=10.00KX Signal A = SE2

Date: 22 Mar 2002EHT = 10.00 kV WD = 4 mm

1mmMag = 20.00KX EHT = 10.00 kV

WD = 6 mmSignal A = SE2Date: 22 Mar 2002

(c) (d)

13.1 Crystal structure and morphology of some previously usedceramic interface phases: (a) h-BN; (b) Ti3SiC2 (courtesy of Y.M. Luo,Tsinghua University); (c) and (d) LaPO4.


garnet (YAG) fibers have also yielded useful information on the design ofthe interfaces.

Sudheendra et al. [5] have carried out a systematic investigation of theparticulate composites of Al2O3 and TiO2 with a variety of monazites such asLnPO4 (Ln = La, Pr, Nd or Gd) by means of scanning electron microscopyin order to check the chemical compatibility of LnPO4 with alumina andtitanium oxide. The studies reveal that all the monazites are compatible withthe oxide ceramics and have favorable mechanical properties in terms ofdamage tolerance and increase in fracture toughness.

13.3.1 Experimental procedure

The nanocrystalline LaPO4 powders were synthesized by mixing phosphoricacid with lanthanum oxide in a water bath [6]. The composite powders ofdifferent LaPO4 composition were ball-milled in ethyl alcohol with agateballs for 24 hours and then dried, and sieved using 100-mesh sift. Details ofmaterial preparation have been reported elsewhere [7, 8]. XRD was carriedout using an X-ray diffractometer (Cu-Ka, Model Automated D/Max B,Japan). Hardness values were determined on polished samples using a 50 Nload in a microhardness tester fitted with a Vickers square pyramidal indenter.Five indents were made on each sample. The hardness was deduced from thediagonal width (2a) of the indentation and the contact load (P) by the followingequation:

HP

a =

1854.4(2 )2 (13.1)

The flexural strength was measured by the three-point bending method withspecimen dimensions of 36 ¥ 4 ¥ 3 mm3, a bending span of 30 mm, and acrosshead speed of 0.5 mm/min at room temperature. The tensile surface ofthe sample was polished with diamond paste down to 0.5 mm and the longedges of the tensile surface were rounded. The elastic modulus was measuredby the load–displacement curve on specimens 36 ¥ 4 ¥ 3 mm3 using an A-2000 Shimadzu universal materials testing machine with a crosshead speedof 0.05mm/min. Fracture surfaces of the composites were observed by scanningelectron microscopy, using SEM LEO 1530, Germany. Samples wereevaporatively coated with carbon or gold to obtain conductivity withoutaffecting observed surface morphology.


Figure 13.2 shows the influence of LaPO4 content on the bending strength ofAl2O3/LaPO4 composites prepared by different sintering routes (pressurelesssintering, hot pressing and spark plasma sintering). It is obvious that


incorporating LaPO4 particles remarkably reduced the fracture strength ofthe composites. From the result, for hot-pressed samples, the monolithicAl2O3 sintered at 1450∞C exhibited a maximum strength of 550 ± 32 MPa,while the strength of 40 wt% Al2O3/LaPO4 composite and LaPO4 ceramicwere 331 ± 41 MPa and 137 ± 18 MPa, respectively. Davis and Marshall [4]noted that the Al2O3/LaPO4 composites are completely unreactive at hightemperature (below their eutectic temperature) and bond poorly to one another.An added benefit was the ability of activating dislocation slip systems inmonazite (LaPO4) under critical shear stresses. Due to the weak bondingbetween Al2O3 and LaPO4 and the activating dislocation slip systems inmonazite (LaPO4), it was evident that microstructure greatly affected themechanical properties.

Figure 13.3 shows the variation of the elastic modulus of the compositeswith increasing LaPO4 addition done by two different sintering routes. Thedecrease in elastic modulus attributed to LaPO4 concentration can be explainedby the rule-of-mixture.

The most important advantage of Al2O3/LaPO4 composites is that theypossess excellent machinability. Hardness is an important parameter as anindication of ceramic machinability. Generally, lower hardness means bettermachinability. Figure 13.4 shows the effect of LaPO4 content on the hardnessof Al2O3/LaPO4 composites made by three sintering routes. Hardness droppedrapidly with increased LaPO4 content. Large amounts of LaPO4 may producea relatively high volume fraction of the grain-boundary phase, which has a

Pressureless sintering at 1600∞CHot pressing at 1450∞CSpark plasma sintering at 1350∞C

0 20 40 60 80 100Content of LaPO4 (wt%)

600

550

500

450

400

350

300

250

200

150

100

Ben

din

g s

tren

gth

(M

Pa)

13.2 Effect of LaPO4 content on the bending strength of Al2O3/LaPO4composites.


Pressureless sintering at 1600∞CHot pressing at 1450∞C


Ela

stic

mo

du

lus

(GP

a)

340

320

300

280

260

240

220

200

180

160

140

120

100

80

13.3 Effect of LaPO4 content on the elastic modulus of Al2O3/LaPO4composites.

Pressureless sintering at 1600∞CHot pressing at 1450∞CSpark plasma sintering at 1350∞C


18

17

16

15

14

13

12

11

10

9

8

7

6

5

43

Vic

kers

har

dn

ess

(GP

a)

13.4 Effect of LaPO4 content on the Vickers hardness of Al2O3/LaPO4composites.


lower hardness than Al2O3 grains. All of the sample density is higher than99% theoretical density, so the decrease of hardness is not due to the lowdensification caused by addition of LaPO4. The resistance of a material tothe formation of a permanent surface impression by an indenter is termedhardness. The formation of weak interfacial bonding between Al2O3 andLaPO4 phases decreases the resistance of the composite to the plasticdeformation. In this system, the crack can propagate along the weak bondinginterface of Al2O3 and LaPO4 and activate the dislocation slip system of theLaPO4 phase [9]. Therefore, the improved crack deflection ability maycontribute to the machinability improvement in the composites. The Vickershardness of the 40 wt% Al2O3/LaPO4 composite and the LaPO4 ceramic byhot pressing was 4.69 ± 0.01 GPa and 4.48 ± 0.18 GPa, respectively. Thesevalues match the requirement of engineering machining using cementedcarbide tools. The indicated LaPO4 ceramic and 40 wt% Al2O3/LaPO4

composite could possess excellent machinability, which mainly arises fromthe ease of chipping during cutting because of weak interfacial bondingbetween Al2O3 and LaPO4.

13.3.3 Microstructure

Figure 13.5 shows the microstructure evolution of the composites fabricatedby pressureless sintering. It is evident that the addition of LaPO4 in compositesretards the grain growth of Al2O3. The back-scattered electron images alsoshow even distribution of the LaPO4 phase in the Al2O3 matrix. Figure 13.6shows the fracture surface of the 40 wt% Al2O3/LaPO4 composites preparedby the same sintering method. A layered LaPO4 phase surrounding the Al2O3

grains is the main fracture surface feature for these composites. It is worthnoting from the SEM images that the secondary phases are primarily locatedat the grain boundaries. Since the thermal expansion coefficient of Al2O3

(~9 ¥ 10–6/∞C) is lower than that of LaPO4 (~10 ¥ 10–6/∞C), a residual tensilestress would be generated at the interface between Al2O3 and LaPO4 duringcooling down after sintering. The existence of such a stress may weaken theinterfacial bonding, leading to regional particle pullout during machining.

13.3.4 Machinability

Using cemented carbide drills, it is possible to successfully machine 40 wt%Al2O3/LaPO4 specimens fabricated by pressureless sintering and hot pressing.Normally, due to their high hardness, diamond tools are used for the machiningof advanced ceramics. However, pure Al2O3 cannot be machined usingcemented carbide drills. As stated above, the layered LaPO4 phase and theweak interface at the Al2O3/LaPO4 grain boundaries are the main reason forthe improvement of the machinability. Both enhance crack deflection and


10 mm 10 mm(a) 10 wt% LaPO4/Al2O3

10 mm 10 mm(b) 20 wt% LaPO4/Al2O3

10 mm 10 mm(c) 30 wt% LaPO4/Al2O3

(d) 40 wt% LaPO4/Al2O3

13.5 Surface morphology of Al2O3/LaPO4 composites by pressurelesssintering.

10 mm10 mm


help avoid catastrophic failure of the material during drilling. Figure 13.7shows the effect of LaPO4 on the drilling rate and mass loss of Al2O3/LaPO4

composites. It is obvious that the machinability of the composites can beimproved by increasing the LaPO4 content. Other important factors stillunder study include the wear on the tool, the cutting force vs drilling rate, thesurface roughness of the workpiece, and the diffusion mechanism betweenthe tool and the workpiece.

The reduced hardness and improved machinability are attributed primarilyto the crack deflection process. It can be seen in Fig. 13.8 that the compositeshowed obvious particle pullout and significant crack deflection alonginterphase boundaries due to the weak interface bonding. The crack deflectionmechanism (absorbing fracture energy and blunting crack tip) could lead toan increase in machinability. As described above, the thermal expansion

2mm

(a)

(b)2mm

13.6 SEM images of fracture surface for 40 wt% Al2O3/LaPO4composite by pressureless sintering.


mismatch between Al2O3 and the second LaPO4 phase can induce residualtensile stresses at the interface between Al2O3 and LaPO4 upon coolingdown from sintering temperatures. This may result in microcracking to dissipatethe strain energy and shield the main crack, thus avoiding catastrophic failureand leading to an increase in machinability.

13.4 Si3N4/h-BN

Silicon nitride (Si3N4) has great promise for engineering applications due toits excellent thermomechanical properties such as high bending strength,high thermal shock resistance and high fracture toughness [10]. However, itis quite difficult to machine after sintering. This disadvantage leads to highmachining cost, and limits Si3N4 applications greatly. A variety of methodshave been utilized to improve the machinability of Si3N4 ceramics, includingporous Si3N4 ceramic [11], microstructure design [12–14], Si3N4/h-BNnanocomposites [15], etc. Reinforcement by nanoparticles is one of themethods to obtain good mechanical reliability. Kusunose et al. [15] reportedthat the Si3N4/h-BN nanocomposite materials possess good machinability.

13.4.1 Experimental procedure

Starting powders (Si3N4, 0.7–1.2 mm, Founder Co., China; h-BN, 0.53 mm,GY Powder Industries, China) were used as raw materials. According to the

0 20 40 60 80 100LaPO4 content (wt%)

Drilling rate (mm/s)Mass loss (g)

Drilling time: 30 secondDrilling force: 49 N

0.25

0.20

0.15

0.10

0.05

0

Dri

llin

g r

ate

(mm

/s)

0.24

0.22

0.20

0.18

0.16

0.14

0.12

0.10

0.08

0.06

0.04

0.02

0

Mass lo

ss (g)

13.7 Effect of LaPO4 content on the drilling rate and mass loss ofAl2O3/LaPO4 composites.


different proportions of Si3N4 and h-BN with 8wt%Y2O3–2wt%Al2O3 (Y2O3–Al2O3, >99.5%, GRINM, China) as sintering aids, the mixtures were carefullyweighed, mixed and milled in a roller mill with ethanol for 48 h, then driedat 80∞C and sieved through a 100 mesh sieve. More details of materialpreparation and characterization have been reported elsewhere [13, 14].

1mmMag = 20.00 KX

EHT = 10.00 kV WD = 10 mm

Signal A = SE2Date: 8 Oct 2001

(a)

1mmMag=20.00KX

EHT = 10.00 kV WD = 10 mm

Signal A = SE2Date: 8 Oct 2001

(b)

13.8 Fracture surface of 30 wt% and 40 wt% Al2O3/LaPO4 compositesby hot pressing.



As is known, h-BN, it is difficult to completely densify at the sinteringtemperature that Si3N4 reaches its theoretical density, due to the chemicalinertness and plate-like structure. Figure 13.9 shows the density change ofSi3N4/h-BN composites with different h-BN concentrations. The reductionin bulk density with increased h-BN content is caused by two factors: thelow density of h-BN (rh–BN = 2.27 g/cm3; rSi N3 4 =3.19 g/cm3), and the low-chemically active nature of h-BN in Si3N4/h-BN composites. Therefore, h-BN reduced the sinterability of Si3N4/h-BN composites, resulting in decreasein bulk density.

The effect of h-BN content on Vickers hardness, flexural strength, andelastic modulus of Si3N4/h-BN composites was investigated. Figure 13.10shows the relationship between composite hardness and h-BN content. Thehardness of Si3N4 with 25% h-BN volume content is as high as HV = 5.67GPa, which matches the request of engineering machining. The hardness isclose to that of machinable mica glass–ceramic (3 GPa) and layered ternarycompounds Ti3SiC2

(4~5 GPa). The Vickers hardness of Si3N4/h-BN compositedecreased with increasing h-BN volume fraction. Easy cleavage of basalplane h-BN platelets causes hardness and fracture strength to decrease withh-BN addition as shown in Fig. 13.11. Along with the addition of h-BN, asteep decrease in the hardness of Si3N4/h-BN composites occurs due to theformation of the weak interface, leading to good machinability. Of course,an increase in porosity will also cause a decrease in hardness. Sometimes,

Theoretical valuesMeasured values

0 5 10 15 20 25h-BN content (vol%)

3.20

3.15

3.10

3.05

3.00

2.95

2.90

2.85

Den

sity

(g

/cm

3 )

13.9 Effect of the h-BN content on the density of Si3N4/h-BNcomposites.


Si3N4/h-BN1750∞C (2h)

0 5 10 15 20 25h-BN content (vol%)

20

18

16

14

12

10

8

6

4

Hv

(GP

a)

13.10 Effect of the h-BN content on the hardness of Si3N4/h-BNcomposites.

Si3N4/h-BN1750∞C (2h)

0 5 10 15 20 25h-BN content (vol%)

750

700

650

600

550

500

Ben

din

g s

tren

gth

(M

Pa)

13.11 Effect of the h-BN content on the bending strength of Si3N4/h-BN composites.


this is more important than a weak interface. However, another 25%h-BN/Si3N4 sample sintered by spark plasma sintering still possesses excellentmachinability even with high densification (99.6% theoretical density). So,when the addition of the weak phase is above 20%, it seems that the weakbonding has a greater effect than increased porosity in the sample, anddominates the improved machinability. As a weak phase, the addition of h-BN will benefit the crack deflection and reduce the composite hardness. Thereduction of composite sinterability also leads to decrease of flexure strengthand increased porosity. It has long been reported that the bending strength ofceramics is governed by porosity, and the same result was observed in thisstudy. An important reason is that Si3N4/h-BN is hard to densify due to theh-BN addition, which produces many pores in the microstructure. The SEMimage of Si3N4/h-BN graded composites proved that the amount of pores inthe specimen increases and the size of pore enlarges with increasing h-BNcontent [14]. At the same time, the pore orientation parallel to the HP directionis similar to those observed by other authors [16]. Figure 13.12 shows thedependence of elastic modulus of Si3N4/h-BN ceramic composites on h-BNcontent. From the result, high h-BN content resulted in reduction of elasticmodulus of Si3N4/h-BN ceramic composites.

13.4.3 Microstructure and machinability

Figure 13.13 shows the microstructures of sintered Si3N4 with and withoutthe addition of 5, 10, 15, 20 and 25 vol% h-BN. SEM micrographs of the

Si3N4/h-BN1750∞C (2h)

0 5 10 15 20 25h-BN content (vol%)

260

250

240

230

220

210

200

190

180

170

Ela

stic

mo

du

lus

(GP

a)

13.12 Effect of the h-BN content on the elastic modulus of Si3N4/h-BNcomposites.


(a) Si3N4

6 mm(b) 5 vol% h-BN/Si3N4

6 mm

6 mm 6 mm(c) 10 vol% h-BN/Si3N4 (d) 15 vol% h-BN/Si3N4

6 mm 6 mm(e) 20 vol% h-BN/Si3N4 (f) 25 vol% h-BN/Si3N4

13.13 SEM micrographs of Si3N4/h-BN composites with differentaddition of h-BN, showing: (a) Si3N4; (b) 5 vol% h-BN/Si3N4;(c) 10 vol% h-BN/Si3N4; (d) 15 vol% h-BN/Si3N4; (e) 20 vol% h-BN/Si3N4; (f) 25 vol% h-BN/Si3N4. Æ ¨ direction corresponds to thehot-pressing direction.


fracture surface perpendicular to the pressing direction confirm the preferredorientation of h-BN plates in the hot-pressed h-BN composites. A similarpreferred orientation of h-BN grains in hot-pressed BN/Al2O3 [17], SiC/BN[18], BN/oxide ceramic [19], and BN/B2O3 has previously been observed.Figure 13.13 also shows that the microstructures of specimens of different h-BN are different in comparison to that of monolithic Si3N4 ceramic. Elongatedb-Si3N4 grains were observed in monolithic Si3N4 and 5 vol%Si3N4/h-BNcomposite, but were not observed in other samples. h-BN grew anisotropicallyin a plate-like configuration. The addition of h-BN retarded the growth ofelongated b-Si3N4 grains as seen in the SEM photographs. A scanning electronmicrograph of the resulting microstructure clearly shows the presence ofacicular b-Si3N4 grains with high aspect ratios compared to the monolithicSi3N4.

Every composite microstructure shows that the layered crystal structureand the preferential orientation of the flaked h-BN grains are perpendicularto the hot-pressing direction, which may be due to the rotation of the h-BNflakes during the viscous flow of the glass phase under hot pressing. h-BNhas a layered structure similar to graphite, with strong bonding within eachlayer and weak bonding between the layers. When a crack tip meets an h-BNgrain, it will propagate either along the interface between the Si3N4 and h-BN grains or along the interlayer within the h-BN grains. In contrast, thecrack deflection and propagation across an h-BN flake are difficult becauseof the strong bonding within each layer of the h-BN grains. The fracturesurface of the composite shows quite a large area of crack deflection, pulloutand cleavage cracking of the h-BN grains. All of these explain the improvementin machinability of Si3N4/h-BN composites.

The mechanism of materials removal of 25 vol% Si3N4/h-BN compositeduring drilling using a cemented carbide drill seems to rely on the cleavageof layered h-BN crystals. During drilling, the cleavage of h-BN crystalslocalizes the fracture of composite at a microscopic scale and allows powderedchips to form easily, giving rise to good machinability. The intercrystalporosity can provide another reason for the improvement of machinabilityby preventing growth of cracks associated with machining. Moreover, itappears that the more porous a ceramic, the weaker, softer and more machinableit becomes. According to Fig. 13.9, the relative density decreased and porosityincreased with the increase of h-BN content, which also resulted in the goodmachinability of Si3N4/h-BN composites.

Figure 13.14 shows the TEM images of fracture surfaces for the 25 vol%Si3N4/h-BN composite. One can see that the plate-like h-BN grains were notwetted by the glass phase and had no obvious strong bonding with the Si3N4

matrix. In some areas, one can even see layered h-BN showing a bendingdirection perpendicular to the c-axis.


13.5 Machinable nanocomposites

As shown above, the improved machinability with increasing concentrationof a second weak phase normally will cause a decrease in fracture strength,except for the Al2O3/Ti3SiC2 system (here weak means that the hardness ofTi3SiC2 is less than that of Al2O3) [20]. How to combine high strength withgood machinability is an inevitable consideration for the wide application ofadvanced ceramics, especially for structural ceramics. In recent years, manyattempts have been made to develop strong machinable ceramics. Amongthem, a machinable nanocomposite is one of the most promising routes [21,22].

(a)

h-BN

150 nm

Glass phase

Si3N4

150 nm

h-BN

C axis (002)

(b)

13.14 TEM observation of machinable 25 vol% Si3N4/h-BN composite.


Because nanocomposites are made from different phases with differentthermal expansion coefficients and elastic moduli, they inevitably developresidual thermal stress during cooling after sintering. Assuming the dispersionphase is spherical particulate in the matrix material, residual stresses can bedeveloped due to differences in the thermal expansion and elastic constantsbetween the matrix and the particles [23]:

s a a D

n nf

m f m

m fm

f

= ( – )

(1 + ) + (1 + )

E TEE

(13.2)

s st

f = –2

(13.3)

where sf is the radial stress and st is the tangential stress; am and af, Em andEf, and nm and nf are thermal expansion coefficients, elastic moduli andPoisson’s ratios of the matrix and the dispersion phase, respectively; and DTis the temperature range during cooling down.

Due to the stress concentration, crack propagation may be pinned by thenanoparticles near the crack tip. The schematic diagrams and TEM images inFig. 13.15 show crack pinning by a nanoparticle and transgranular fractureinduced by an intragranular particle. These experiments also show that crackpinning can give rise to nanoparticle pullout. When crack growth is pinnedby a nanoparticle, the crack can penetrate through the nanoparticle by breakingit or can propagate along the nanoparticle and matrix. But due to the residualstress, the crack dominates over propagation along the phase boundary. Themain mechanisms involve deflection of cracks or secondary cracking causedby the presence of a weak layer, weak interfaces, residual stress, or othermicrostructural defects. If the crack propagation is along the weak interfacebetween the matrix and the soft phase, the main materials removal mechanismis grain pullout. This will avoid catastrophic fracture during machining. Theweak bonding and weak phases are the main sources of energy dissipationand damage tolerance. In Al2O3/LaPO4, the interfacial toughness is sufficientlylow to satisfy the criterion of He and Hutchinson for a normally incidentcrack to deflect along the interface rather than cross it. This system is stablefor long periods in air at temperatures at least as high as 1600∞C. Dislocationsare the only line defect in a solid material. As we known, the movement ofdislocations is the elementary mechanism of plastic deformation of manycrystalline materials, especially in ceramics. A more detailed investigation isstill needed to understand the relation between the improved machinabilityand formation of dislocations induced by residual thermal stress.

13.6 Conclusions

Design and fabrication of strong machinable ceramics is emerging as aninevitable requirement for flexible use of advanced ceramics, especially for


structural ceramics. However, the extremely high hardness of ceramics makesconventional machining very difficult or even impossible. In the past 10years, many researchers have focused on the improvement of ceramicmachinability. Normally, compound design of machinable ceramics is usedto improve the machinability of ceramic materials. This method introducesa weak interphase or layered structure material into the matrix to facilitatecrack deflection and propagation during machining and therefore avoid brittlefracture. However, the improved machinability with increasing concentrationof a second weak phase normally will cause a decrease in fracture strength.At present, machinable nanocomposite design is considered one of thepromising methods to combine high strength with good machinability. More

13.15 Schematic and TEM image of crack pinning by a nanoparticleand transgranular fracture induced by intragranular particle.

Crack

Al2O3 matrixLaPO4 particle

100 nm(a) 100 nm(b)

(c)100 nm


investigation is necessary on processing procedures, sintering thermodynamics,residual thermal stress and microstructure details to understand the real reasonfor the improved machinability.

13.7 References

1. Niihara, K., New design concept of structural ceramics – ceramic nanocomposites,J. Ceram. Soc. Jap., 99(10): 974–982 (1991).

2. Parthasarathy, T.A., Boakye, E., Cinibulk, M.K., and Perry, M.D., Fabrication andtesting of oxide/oxide microcomposites with monazite and hibonite as interlayers, J.Am. Ceram. Soc., 82(12): 3575–3583 (1999).

3. He, M.Y., and Hutchinson, J.W., Crack deflection at an interface between dissimilarelastic materials, Int. J. Solids Structures, 25(9): 1055–1067 (1989).

4. Davis, J.B., and Marshall, D.B., Machinable ceramics containing rare-earth phosphates,J. Am. Ceram. Soc., 81(8): 2169–2175 (1998).

5. Sudheendra, L., Renganathan, M.K., and Raju, A.R., Bonding of monazite to Al2O3

and TiO2 ceramics, Mater. Sci. Eng. A, 281(1–2): 259–262 (2000).6. Wang, R.G., Pan, W., Chen, J., Fang, M.H., Cao, Z.Z., and Luo, Y.M., Synthesis and

sintering of LaPO4 powder and its application, Mater. Chem. Phys., 79(1): 30–36(2003).

7. Wang, R.G., Pan, W., Chen, J., Jiang, M.N., Luo, Y.M., and Fang, M.H., Propertiesand microstructure of machinable Al2O3/LaPO4 ceramic composites, CeramicInternational, 29(1): 19–25 (2003).

8. Wang, R.G., Pan, W., Chen, J., Fang, M.H., Jiang, M.N., and Cao, Z.Z., Microstructureand mechanical properties of machinable Al2O3/LaPO4 composites by hot pressing,Ceramic International, 29(1): 83–89 (2003).

9. Davis, J.B., et al., Influence of interfacial roughness on fiber sliding in oxide compositeswith La-monazite interphases, J. Am. Ceram. Soc., 86(2): 305–316 (2003).

10. Rouxel, T., High temperature mechanical behavior of silicon nitride ceramics, J.Ceram. Soc. Jap., 109(6): 89–98 (2001).

11. Kawai, C., and Yamakawa, A., Machinability of high-strength porous silicon nitrideceramic, J. Ceram. Soc. Jap., 106(11): 1135–1137 (1998).

12. Kawai, C., and Yamakawa, A., Effect of porosity and microstructure on the strengthof Si3N4: designed microstructure for high strength, high thermal shock resistance,and facile machining, J. Am. Ceram. Soc., 80: 2705–2708 (1997).

13. Wang, R.G., Pan, W., Chen, J., Jiang, M.N., and Fang, M.H., Fabrication andcharacterization of machinable Si3N4/h-BN functionally graded materials, Mater.Res. Bull., 37(7): 1269–1277 (2002).

14. Wang, R.G., Pan, W., Jiang, M.N., Chen, J., and Luo, Y.M., Investigation of thephysical and mechanical properties of hot-pressed machinable Si3N4/h-BN compositesand FGM, Mater. Sci. Eng. B, 90(3): 261–268 (2002).

15. Kusunose, T., et al., in Innovative Processing and Synthesis: Ceramic, Glass andComposite, edited by Bansal, N.P., Logan, K.V., and Singh, J.P., Westerville, OH:American Ceramic Society, 1997, pp. 443–454.

16. Oliveira, F.J., Carrapichano, J.M., Silva, R.F., and Vieira, J.M., Sintering of Si3N4-BN composites, Sil. Ind., 63(5–6): 69–72 (1992).

17. Goeuriot-Launay, D., Brayet, G., and Thevenot, F., Boron nitride effect on the thermalshock resistance of an alumina based ceramic composite, J. Mater. Sci. Lett., 5: 940–942 (1986).


18. Ruh, R., Bentsen, L.D., and Hasselman, D.P.H., Thermal diffusion anisotropy ofSiC/BN composites, J. Am. Ceram. Soc., 67: C-83 (1984).

19. Trice, R.W., and Halloran, J.W., Investigation of the physical and mechanical propertiesof boron nitride/oxide ceramic composites, J. Am. Ceram. Soc., 82(9): 2563–2565(1999).

20. Luo, Y.M., Li, S.Q., Chen, J., Wang, R.G., Li, J.Q., and Pan, W., Effect of compositionon properties of alumina/titanium silicon carbide composites, J. Am. Ceram. Soc.,85(12): 3099–3101 (2002).

21. Kusunose, T., Sekino, T., Choa, Y.H. and Niihara, K., Machinability of silicon nitride/boron nitride nanocomposites, J. Am. Ceram Soc., 85(11): 2689–2695 (2002).

22. Wang, X.D., Qiao, G.J., and Jin, Z.H., Fabrication of machinable silicon carbide–boron nitride ceramic nanocomposites, J. Am. Ceram Soc., 87(4): 565–570 (2004).

23. Li, R.T., PhD dissertation, Tsinghua University, Beijing, 2002.

Part IV

Refractory and speciality ceramiccomposites

357

359

14.1 Introduction

The spinels are a class of double oxide of general formula AB2O4: industriallyimportant members of this class include aluminates (e.g. MgAl2O4), ferrites(e.g. MgFe2O4) and chromites (e.g. MgCr2O4).

Magnesium aluminate spinel (MgAl2O4) is an important constituent ofmagnesia-based refractory materials. The melting point of MgAl2O4 is 2135∞C.There are no natural deposits of MgAl2O4, which is therefore normally obtainedby reaction of mixtures of magnesium and aluminium oxides. The theoreticalstoichiometric composition of magnesium aluminate spinel is 71.68% Al2O3

and 28.32% MgO by weight, but compositions can vary because of a limitedrange of solid solution. Its density is 3.579 Mg m–3, approximately the sameas MgO (3.583 Mg m–3).

Commercial sintered magnesia–spinel refractory materials are dividedinto three categories: magnesia rich, stoichiometric, and alumina rich.1 Typicalproperties of magnesia rich spinel bricks are given in Table 14.1.2

Spinel is used as an additive in magnesia rich bricks, for example forcement kiln linings. Magnesia rich spinel bricks are used in the cooling zoneand in the upper side of the sintering zone of the cement kiln.3,4 In addition,spinel particles are added in various proportions to MgO in order to improveits thermal shock resistance. Magnesia–spinel materials give significantly(two to three times) longer service lives than other basic bricks such asconventional magnesia chrome.2 The reason for the improvement in thermalshock resistance is linked to the large difference in thermal expansioncoefficient5 between magnesium oxide (mean value ~13.5 MK–1) and spinel(~7.6 MK–1). This difference leads to the development of large tensile hoopstresses, and eventually extensive microcracking, around spinel grains duringcooling from fabrication temperatures in excess of 1600∞C. Thermal shockis particularly severe during the kiln heating and cooling periods, whentemperature changes can be large and rapid, leading to significant thermal

14Magnesia–spinel (MgAl2O4) refractory

ceramic composites

C A K S E L, Anadolu University, Turkey andF L R I L E Y, University of Leeds, UK


stresses. Some transient healing of the cracks may occur on subsequentreheating during use of the refractory.1,6–8

The stoichiometric type is similarly used in cement kiln linings, as well asan alumina-based castable in ladles, which can also have good resistance tocorrosion and erosion.9 The alumina rich type is being studied for use inalumina-based castables in order to improve resistance to slag* penetration;85–90 wt% Al2O3 is suggested as the optimum. A higher alumina contentincreases the spinel bonding and improves strength, and increased spinelbonding results in better spalling resistance.10 Alumina rich magnesia–spinelis also used in steel and aluminium production, petrochemical applications,and glass melting furnaces.3

A major reason for the use of MgAl2O4 compared to other spinels (suchas magnesia–chrome) is its better resistance to thermal shock and alkaliattack2 (Fig. 14.1). Magnesia–chrome refractories were used for many yearsas high strength hot-face refractories in a range of systems, including rotarycement kilns and steel-making vessels,11 but there is risk of the contaminationof ground water by hexavalent (VI) chromium ions leached from wastematerials. Hence, a second significant reason for the escalating interest inMgAl2O4 is to avoid the use of refractories containing chromium oxide.Magnesium aluminate spinel is therefore being used as an alternative secondphase in magnesia–spinel and alumina–spinel refractories, allowing a wide

Table 14.1 Typical properties of magnesia rich spinel bricks2

Properties Magnesia–spinel Magnesia–spinelbrick with brick with sinteredalumina spinel

MgO (%) 96–86 90–80Al2O3 (%) 3–8 9–18Fe2O3 (%) 0.1–4 0.1–0.5CaO (%) 0.4–2 0.4–1.0SiO2 (%) 0.2–4 0.2–0.6Bulk density (Mg m–3) 2.85–2.95 2.90–3.00Apparent porosity (%) 16–20 14–18Cold crushing strength 30–90 40–80

(N mm–2)Refractoriness under load (∞C) 1550–1700 >1700Thermal shock resistance 40–100 >100

950∞C/airThermal conductivity

(W m–1 K–1)500∞C 3.0–4.0 3.0–4.01000∞C 2.4–3.0 2.8–3.7

*Slag is the product of reaction between fluxing materials and a refractory furnacelining116

Magnesia–spinel (MgAl2O4) refractory ceramic composites 361

range of compositions and types for a large number of applications to beproduced.9,12,13 The environmental hazard posed by the conversion of insolubletrivalent (III) chromium to the soluble hexavalent state allows Cr(VI) ionleaching from waste magnesite–chrome refractories. Cr(VI) has been associatedwith allergic skin ulceration and carcinomas in humans.11 When chrome orereacts with alkalis to form potassium chromate or potassium dichromate(containing Cr(VI)), this can result in the destruction of the brick.5,14 Hexavalentchromium diffuses from the refractory into the cement clinker, and increasesthe risk of toxic reactions during processing of the cement.2

In cement kilns, low viscosity clinker can also react with magnesia–chromespinel to form relatively low melting point compounds, especially monticellite(CaO.MgO.SiO2).15,16 There are eutectic points at 1360 and 1490∞C in themonticellite region.17 In the CaO–Al2O3 binary system near the compositionCaO.6Al2O3, there is a eutectic point at 1395∞C.18 The spinel content musttherefore be kept to a minimum, and it is necessary to establish the optimumamount. One widely used type of refractory is based on material consistingof coarse (mm) grain size magnesite, bonded with a (mm) grain-size magnesite–spinel phase.

The main advantages6,7,13,19-22 of using magnesia–spinel bricks in cementkilns can be summarised as:

∑ Low thermal expansion coefficient of magnesia–spinel bricks∑ High resistance to thermomechanical stress

= very good = good = average = low

Resistance to Dolomite Magnesia–chromite Magnesia–spinel

Reducing atmosphere

Free SO2/O3

CO2

Free K2O/Na2O

Clinker liquids

K2SO4

KCl

Thermal shocks

Stress (kiln shell)

Abrasion

14.1 Comparison of resistance of dolomite, magnesia–chromite andmagnesia–spinel to environmental attack.2


∑ Chemical resistance to oil and ash deposits∑ High resistance to corrosion and changes in kiln atmosphere∑ Low content of secondary oxides which results in minimal alteration in

structure of the hot face in service∑ Elimination of chromite that makes the brick less susceptible to alkali

attack in service∑ No toxic Cr(VI) ions leached from waste materials∑ White cements can be made without discoloration problems caused by

transition metal cations.

14.2 Crystal structures

MgO has the NaCl crystal structure;23,24 each magnesium ion is coordinatedby six O2– ions and each O2– by six Mg2+. The structure of a-Al2O3 consistsof close-packed sheets of O2– ions stacked in the sequence A–B–A–B, forminga hexagonal close-packed array of anions.25 The cations are located in two-thirds of the octahedral sites. The structure of a-Al2O3 results in coordinationnumbers of 6 and 4 for the cation and the anion, respectively.25

Spinels of the general formula AB2O4, such as magnesium aluminatespinel, MgAl2O4, are based on a face-centred cubic (fcc) array of oxygenions.26 There are two types of spinel: in normal spinels the A2+ ions are ontetrahedral sites and the B3+ ions are on octahedral sites (MgAl2O4); ininverse spinels, the A2+ ions and half the B3+ ions are on octahedral sites,with the other half of the B3+ ions on tetrahedral sites, B(AB)O4: an exampleis FeMgFeO4.

26 There are two kinds of interstice in a close-packed lattice:27

octahedral sites, defined by six ions, and tetrahedral sites, defined by fourions. In MgAl2O4, the oxygen ions are in a face-centred cubic, or cubic closepacked, array. A unit cell contains eight Mg2+ ions, 16 Al3+ and 32 O2– ions;there are 32 octahedral interstices and 64 tetrahedral interstices.26 In thenormal spinel, the octahedral sites are occupied by the smaller trivalent ions,and the tetrahedral sites are occupied by the larger divalent ions. There aretwice as many filled octahedral sites as tetrahedral sites, corresponding tothe numbers of trivalent and divalent ions in MgAl2O4: one-eighth of thetetrahedral and half of the octahedral sites are filled.26

14.3 Production of MgAl2O4

Commercially synthetic spinel is produced in four main district forms:12,28

∑ Electrofused spinel∑ Sintered spinel clinker*

*Clinker is a material that has been fully pre-reacted at high temperature in granulatedform.116


∑ Sintered magnesia–spinel clinker∑ Calcined spinel.

In the fusion process calcined Bayer Al2O3 and MgO powders are mixedin stoichiometric ratio, and melted in an electric arc furnace at 2200∞C. Aftercooling and milling a dense spinel is obtained.

Sintered spinel clinker has been developed as a more economical rawmaterial than fused spinel, which it replaces in many applications. Al2O3,MgO and a selected binder (organic polymer or inorganic) are pulverisedand mixed in a tube mill. After forming and drying, the mixture is sinteredat temperatures above 1600∞C in a rotary kiln until fully dense.

MgO–spinel clinker is obtained by the same procedure, but with a substantialproportion of free MgO (75–85%). MgO–spinel is produced by sintering attemperatures of 1600–1800∞C and a dense product is obtained. MgO–spinelrefractory has better high temperature mechanical properties in arc furnaceroof application compared to magnesia–chrome refractories, but it has lowerstrength than a high-alumina refractory.

High-density spinel refractory brick is made by calcining (1200–1300∞Cfor about 3 h) compacted powders and then sintering at 1700∞C. The extentof spinel formation increases by calcination; however, the powder mixture isnot completely converted to spinel. Typically 10–15% of a–Al2O3 and 5–10% of MgO are observed by XRD, depending on the conditions of temperature,time and particle size of the reactants.

14.4 Densification

Grain size and boundaries, impurities, additives, particle size, porosity, sinteringtemperature and time, heating and cooling rate, and shaping practice play animportant role in controlling many physical, mechanical and chemical propertiesof magnesia-based bricks.29

14.4.1 Sintering of MgO

The main difficulty in sintering of MgO is its chemical instability withrespect to water and atmospheric humidity. Hydration of MgO causes extensiveswelling. Dehydration during drying and in sintering is therefore connectedwith an extremely large shrinkage, which can be associated with the formationof cracks in sintered materials. Precalcination of the MgO powder (especiallysubmicrometre powder) makes the initial material less susceptible to hydration.The use of alcohol rather than water as a mixing or milling medium alsoreduces hydration.30

The sinterability of MgO powder depends on powder purity, particle size,precalcination stage and temperature.31 MgO of purity ranging from 98.5%


to 99.3%, precalcined at 600–1300∞C and sintered at 1650–1870∞C, givesmaterial of density 96–98%.30,32,33 Hot-pressing >50 nm MgO powder at1600∞C and ~15 MPa for 1 h gives material of relative density ~98%.34 A10–50 nm MgO powder containing 0.5–1 mm agglomeration and of purity>99.9% sintered at 1450∞C for 5 h gives 95% dense material of grain size20–30 mm.35 Rapid grain growth leads to the entrapment of pores; pores farfrom the grain boundaries will disappear more slowly than those on theboundaries. For this reason, the location and size of the closed porosity playan important role in controlling the maximum sintered density.

14.4.2 Sintering in the MgO–Al2O3 system

MgO and Al2O3 powders of ~1 mm particle size react to form spinel andsinter, but it is difficult to obtain full sintered density even at temperatures ashigh as 1750∞C. If high density is to be obtained, precalcination at a lowertemperature (1300∞C for 3 h) is essential.32,36,37 This stage allows almost fullreaction to spinel to occur. For 55–79% of pre-reacted spinel, final densitiesof 95–96% can be obtained at 1700–1750∞C in 15–60 min.

The minimum temperature required to obtain a fully dense product increaseswith Al2O3 content from the stoichiometric spinel (71.68 wt% Al2O3) to 85wt% Al2O3, requiring from 1650∞C to ~1680∞C. The sinterability of spinelthus depends markedly on stoichiometry; it then decreases steadily to 1650∞Cwith further increased alumina content (93 wt%).38

14.4.3 Solid state reactions

The formation of spinel from magnesia and a-Al2O3 is associated with avolume expansion of ~8%.39 Solid state reaction between MgO and Al2O3 toform spinel occurs at the Al2O3–MgAl2O4 and MgO–MgAl2O4 interfaces.The reaction between these is by counterdiffusion of the Al3+ and Mg2+ ionsthrough the rigid oxygen lattice or the spinel phase. Three Mg2+ ions diffusefor every two Al3+ ions that diffuse in the opposite direction. Three moles ofspinel are formed at the Al2O3–MgAl2O4 interface for every mole formed atthe MgO–MgAl2O4 interface because of ionic charges andstoichiometry.32,33,40,41 This is illustrated by the following schematic sectionthrough a growing MgAl2O4 layer.

I II

MgO MgAl2O4 Al2O3

Æ 3Mg2+ 2Al3+ ¨At I 2Al3+ + 4MgO Æ MgAl2O4 + 3Mg2+

At II 3Mg2+ + 4Al2O3 Æ 3MgAl2O4 + 2Al3+


14.4.4 Liquid phase sintering in the MgO–spinel system

In the MgO–spinel system, the diffusion rates of Mg2+ and Al3+ are acceleratedby adding components of the liquid-forming system containing CaO, SiO2

and Fe2O3; and densification can be achieved with 2–3 wt% of additive at1600∞C. The rate of grain growth is decreased because these additions restrictdiscontinuous grain growth.42,43 Using CaO and SiO2, the densities obtainedat 1600∞C for ~30 min were equal to those obtained without addition at1800∞C for ~30 min.36

14.4.5 Grain growth in the MgO–spinel system

Approximately 200 nm 99.8% purity spinel powder begins to densify attemperatures above 1100∞C. Neck growth between grains takes place at1100∞C and continues up to ~1500∞C because of the slow changes in grainsize distribution. Grain growth begins to take place at around 1100∞C andincreases significantly from 1500∞C up to 1625∞C; however, the grains stopgrowing on further heating to 1700∞C.44 The average grain size at 1500∞C istwice as large as at 1100∞C. A halt of grain growth above 1625∞C can beassumed to be because the migration of grain boundaries is restricted by thepresence of pores.44

Spinel particles are precipitated on cooling largely at the MgO grainboundaries, but some are trapped within growing MgO grains.28

14.5 In situ formed/preformed spinel-based

refractories

In situ formed spinel-based refractories use the addition of an alumina sourceto form the spinel in situ by reaction with the magnesia matrix during sintering.A strong bond between spinel grains and the magnesia matrix is formed.45

Initial spinel formation occurs around the periphery of the alumina particleand proceeds towards the particle centre.46 The strongly bonded peripheralspinel and a hollow core in these granules is claimed to give better fracturetoughness.46 The origin of the hollow core is not clear. The strength andhigh-temperature fracture characteristics depend strongly on the level ofimpurities in the magnesia, and the type and distribution of secondaryphases.45,47 The complete conversion of the granular alumina is slow duringmanufacture of the spinel. Residual-free alumina will continue to form spinelwith associated expansion in service, at a rate that depends on temperatureand time. A smaller alumina particle size gives faster spinel formation, andthe product magnesia–spinel composite also has better strength.33,45,46,48,49

The second type, which is technologically more advanced, is usuallymade by mixing 10–25% sintered or fused synthetic spinel with 75–90%


sintered magnesia clinker. Sintered magnesia–spinel clinker is formed byheating at ~1700∞C.50 In contrast to the in situ reaction of alumina, whichdevelops a strong spinel bond with magnesia, preformed spinel grain haslittle tendency to bond to the magnesia. Preformed spinel in a magnesiamatrix has a peripheral gap around the spinel grain, which appears to act asa crack initiator, and cracks propagate into the surrounding magnesia matrix.46

Cracks often emanate from sharp corners. Fracture is arrested when thecrack meets the peripheral gap of another spinel granule.46

The properties of MgO–spinel refractories containing preformed spinelare more strongly influenced than the in situ formed spinel refractory bysintering temperature. Fracture toughness increases significantly with increasingsintering temperature; however, the Young’s modulus decreases.46 The high-temperature behaviour of spinel is explained in Section 14.6.4.

14.6 Industrial applications and properties of

magnesia–spinel materials

14.6.1 Type and effect of the additives in magnesia–spinel refractories (CaO, SiO2 and Fe2O3)

Keeping the level of CaO low prevents the formation of low melting pointcalcium aluminate, which can lead to fractures and destruction of the lining,in the magnesia–spinel brick structure during firing.51 A carefully adjustedCaO content (as shown in Table 14.2) forms a protective coating in thesintering zone. A larger amount of silicate in the microstructure formsinterconnected silicate networks, which can act as crack propagation pathsand decrease the resistance to spalling.6 The optimum value of the CaO/SiO2

ratio and reasons are given in Table 14.2.4

Table 14.2 Brick components of basic magnesia–spinel bricks, their most importantreactions with the kiln feed and volatile components from the kiln atmosphere andthe resulting demands on the bricks4

CaO/SiO2 I. C/S >> 1.87 Decrease of(C/S) ratio CaOfree + MgAl2O4 Æ refractoriness,*(wt%) calcium aluminates abrasion†

CaOfree + SO3 Æ CaSO4 Infiltration

II. C/S = 1.87 Formation of CaO/SiO2 in thehighly refractory brick approximatelycompounds; few 1.7–2.2 if S(C+S) > 1.25reactions

III. C/S << 1.87 Decrease of2MgO.SiO2 + 2CaO.SiO2 Æ refractoriness2(CaO.MgO.SiO2)

*Refractoriness is the ability of a material to withstand high temperatures that is evaluated interms of the Pyrometric Cone Equivalent.116

†Abrasion is wear caused by the mechanical action of a solid.116


A reducing/oxidising atmosphere is important for the performance of thebricks, especially in kilns fired with solid fuel, since the reactions seen inTable 14.2 are combined with dramatic changes in volume, which result inthe destruction of the brick where unfired particles can lead to reducingatmosphere. This condition affects the brick severely and may lead to prematurewear. A magnesia–spinel brick for the cement industry should therefore haveas little Fe2O3 as possible (preferably under 1%).52 Service life in the transitionand the cooling zone of the rotary cement kiln compared to other basicbricks is longer. This type of magnesia–spinel brick can also be used in theupper side of the sintering zone of the cement kiln.53

The effect of phase changes from MgO–FeO to MgO–Fe2O3 on the bondedtexture might be disadvantageous in developing high thermal spallingresistance.54 A reducing atmosphere accelerates the decomposition of alkalinesalts, which causes their corrosive power to be increased.55 When MgO–FeOis oxidised, the volumetric expansion is about 8–23%, depending on theamount of FeO. The dissolution of Fe2O3 in MgAl2O4 leads to the formationof cation vacancies, which introduces defects in the host lattice. Duringferrite formation the reaction rate depends on temperature, grain sizedistribution, density and also the partial oxygen pressure because theconcentration gradients over the reaction layer increase with decreasing partialoxygen pressure. The ferrous content of the final product depends on thehomogeneity of the starting mixture.56,57

14.6.2 Applications in rotary cement kilns

A rotary cement kiln can be divided into five zones,2,58–60 which are takensequentially from the beginning (inlet cone) where the raw material isintroduced. Temperatures in this zone are in the 800–1000∞C range and therefractory installed is ordinarily an aluminous firebrick. The calcining zoneuses the low heat transfer properties of thermally insulating MgO–spinelbrick, and high-alumina brick. This is where the corrosive effects of volatilesand alkaline salts manifest themselves. Semi-insulating brick may also beused for lining in this zone, depending upon its mechanical properties. The(upper) transition zone is characterised by mechanical stresses, thermal shockand chemical attack. The chemical attack occurs by reaction between therefractory and the silica in the clinker melt, and also alkali–sulphate interactionswith volatile species in the kiln atmosphere. Unstable clinker coating leadsto thermal spalling of the refractory lining. The sintering zone (the burningzone) is characterised by high-temperature (1500–1800∞C) corrosion andextreme wear on the refractory lining. For this zone magnesia–spinel ormagnesia–chrome bricks, or less commonly magnesia-enriched dolomite,are used. The burning zone of a rotary cement kiln is normally lined withbasic bricks. In this area, the lining quality must be capable of retaining a


stable clinker coating, which is built up on the lining to afford protection, tocreate a barrier between the hostile elements and the refractory to preventerosion. Although the kiln temperature is highest in this region, the thermalgradient is not excessive and the refractories lining this zone are thereforenot subjected to the same degree of thermal shock. Dolomite-based refractoriesare the dominant form of brick used in this zone, due to their ability tomaintain a clinker coating but also for their high degree of refractoriness.Less frequently magnesia–spinel or magnesia–chrome bricks are employed.The cooling zone (lower transition zone) also depends on the abrasion andspalling resistance of MgO–spinel brick. It is believed that conditions aremore severe in the transition zones where the coating is thinner and lessstable, because thermal gradients are higher than in the other zones, resultingin extreme thermal shock damage. The service life in the lower transitionzone is particularly severe due to the higher operating temperatures. Servicelife using the magnesite–spinel refractory in the lower transition zones hasimproved. For example, magnesia–spinel brick lost at least 60% of the originallining thickness in a one-year campaign, however, service life using magnesite–chrome bricks is approximately eight months.61,62

Brick wear in solid-fuel firing is accelerated by falling of the coatingcaused by the instability of the flame. Coating loss is considered to be afrequent occurrence in rotary cement kiln transition zones. When coatingloss occurs, the refractory lining is subject to a very rapid rise in temperature,which can cause thermal shock damage to the refractory lining.61

Accumulation of alkalis from the fuel and reducing atmosphere is causedby unburned carbon from the precalciner. The alkali minerals volatilise inthe high temperature burning zone and are carried by the gas stream to thefeed end of the rotary kiln where they condense and return to the feed. Thiscycle within the cement kiln has a concentrating effect on the alkalis. It iscommon to find considerable concentrations of alkali salts deposited in thebrick pores when examining basic brick taken from cement kiln burningzones. These alkali salts penetrate through the pore structure and freezetowards the cold end of the brick. When chrome is present in the basic brick,it is also possible to form alkali chromates. Alkali attack of magnesite–chrome brick deteriorates the bonding, causing loss of strength. MgO–spinelbricks are better than other basic refractory bricks in solid-fuel and oil-firedcement kilns. The life of MgO–spinel bricks was more than 1.5–2 times thatof direct-bonded magnesia–chrome bricks in each zone, especially in thecooling zone and upper transition zone. The use of MgO–spinel refractorymaterials in the sintering zone of cement kilns is also being considered.61

14.6.3 Problems in the rotary cement kiln

Service conditions, selection and installation of refractories, quality of rawmaterials and refractory materials are important parameters in increasing the


lifetime of refractory materials. Generally, problems in a rotary cement kilncan be classified into three groups:51,63–65

∑ Thermal problems: thermal shock, excessive thermal load, infiltration ofsilicates, migration of silicates

∑ Mechanical problems: thermal expansion, concentric stress cracks,displacements, deformation of the kiln shell, formation of grooves, forcesoriginating from retaining rings

∑ Chemical wear problems: infiltration of alkali salts, redox effects, hydrationcracks, corrosion of chrome ore.

The function of the refractory lining in the rotary kiln is to resist thermalstresses and withstand the abrasive effects of the cement clinker. Thermalstress fracture has been attributed to rapid heating and cooling of the refractorylining component. The temperature gradient during rapid heating or coolingdevelops thermal stresses. The lining reduces maximum tensile stress duringheating. Fracture of the refractory component results in lining deterioration.During cooling of a hot lining, faster contraction is observed on the hot faceof the lining than on the centre of the brick. The maximum tensile stressoccurs at the (cooled) surface, whereas the maximum tensile stress duringheating is in the centre and is half the maximum tensile stresses duringcooling.66 Further cracks are more stable when produced by cooling thanthose produced by heating. Therefore a higher temperature difference isneeded to initiate crack propagation on heating but crack propagation ismore catastrophic.46

Stress begins to build near the hot face at low temperatures. The Young’smodulus of a brick is lower at the maximum hot face temperature of ~1450∞C.Thus the stress is reduced near the hot face. The maximum stress is transferredtowards the interior of the brick as the hot face temperature rises. High stressis developed towards the hot face of the brick. The lower the thermalconductivity, the higher the rate of thermal expansion, and the higher theYoung’s modulus, the greater the stress.67

During the steady state heat flow conditions in cement kiln operation,high levels of thermal stress should not occur. Critical levels of thermalstress for fracture would be associated with heat flow conditions of theheating and cooling periods. Sudden changes in temperature produce thermaltensions contributing to breakage of bricks. Temperature distributions innon-steady-state conditions are mostly noticeable during the initial stage ofheating when temperature differences through the lining are greatest. Thesecan become particularly severe when the initial heating by oil is changed tosolid fuel, which can result in a sudden and marked increase in the heatingrate, sometimes resulting in thermal shock damage and expansion. This problemcan be minimised by the practice of holding the kiln at about 800∞C for a fewhours during warm-up to allow a near-uniform temperature distribution to be


obtained. Coating formation will decrease the temperature distribution throughthe lining, and loss of coating causes thermal shock due to flame impingement.68

Methods used to improve brick lifetime are the production of stable coatingsin the burning zone, the control of ovality of the tyres and kiln alignment,avoidance of frequent kiln stoppages, and slow heating and cooling rates.62,69,70

14.6.4 Performance

Effect of expansion mismatch on Young’s modulus at high temperatures

Fired MgO–spinel brick has high resistance to thermomechanical* stress,which is a result of a combination of direct crystal bonding and the thermalexpansion mismatch between magnesia and spinel grains. Direct bondingprovides high temperature strength. The mismatch is produced by the differencein thermal expansion between the magnesia and the spinel.6,71

As spinel is added to MgO, a thermal expansion mismatch is introducedinto the system and microcracking occurs. This lowers the Young’s modulusand high-temperature strength.72–75 There is a slightly descending trend whenthe temperature rises and a small peak is shown at ~1200∞C, which is theresult of expansion mismatch between the silicate phases and spinel. Thelater decrease in Young’s modulus is the result of softening of silicate phases.On further cooling, the expansion mismatch causes cracks, which result in adecrease in Young’s modulus and strength. On reheating from room temperature,the cracks close up and partial healing can occur.8,72

Thermomechanical behaviour of stoichiometric spinel at hightemperatures

Young’s modulus, fracture toughness and strength of stoichiometric spinelhave been studied from room temperature up to 1200–1500∞C.76–78 In thelow-temperature region, this material behaves elastically, whereas in thehigh-temperature (>800∞C) region, plastic deformation is possible. At lessthan 800∞C, the fracture toughness (KIC) decreases with increasing temperature;however, KIC increases rapidly with increasing temperature above 800∞C.The decrease in KIC for the low-temperature (<800∞C) elastic region is explainedby the decrease of Young’s modulus with increasing temperature. In thehigh-temperature region, strength decreases with temperature, while KIC

increases from 800∞C up to 1200–1500∞C.76–79 This indicates that the fracturebehaviour of spinel at temperatures greater than 800∞C is governed by non-linear elastic processes. Several explanations can be given for transition to

*Thermomechanical properties, e.g. softening under load, creep in compression,refractoriness under load and thermal shock resistance.


totally intergranular fracture over 800∞C: for example, the loss of grain-boundary coherency from the presence of impurities concentrated at thegrain boundaries, and a crack-tip plasticity enhanced crack-propagationmechanism e.g. secondary cracking in the primary crack tip.76,77,80

At the sintering temperature, liquid phases are formed and would remainas glass at room temperature due to high cooling rates. Grain boundarymicrocracking was observed because of dislocation* and grain boundarysliding due to secondary glassy phase softening. Grain boundary microcrackingcan lead to a decrease of strength during fracture, even if toughness increases.The fracture path changes from intergranular at low loading rates totransgranular for high loading rates.79

Hot strength

Hot strength of MgO–spinel is reduced in the presence of commercial spineldue to the low melting point of the impurity phases (CaO.6Al2O3 andmonticellite). As the spinel content increases up to 40%, the hot strengthincreases almost linearly due to the decrease in the impurity content of MgO.However, for spinel amounts over 40%, there was no marked change in thehot strength, even though the impurity content in MgO (>2%) continued todecrease with further addition of spinel up to 80%.28

Slag resistance and corrosion

Spinel-bonded magnesia shows better slagging and spalling resistance dueto the development of strong solid solution bonds between MgO and spinelon firing at 1600–1800∞C. The arc furnace slag resistance improves with thepresence of 10% spinel in the magnesia, but no further improvement occursuntil the spinel content exceeds 40%. The resistance to basic slag deterioratessharply with the addition of more than 40% spinel.28 Magnesia–spinel bricksperform better than basic slags when compared with direct bonded magnesia–chromite and mullite refractories.81

Since the slag is the most corrosive component in the melt, its compositionhas a critical effect on the corrosion mechanism. Corrosion resistance in therotary slag test is evaluated on thickness loss and the depth of penetration.The stoichiometry of the spinel used strongly influences the corrosionbehaviour; for example, alumina-rich spinel addition increases the resistanceto slag penetration and wear. Furthermore, the size of the spinel particlesadded also affects the wear rate: fine particle additions are more effective in

*Dislocations and line defects are relatively easy to move by low stresses, making thecrystals weaker and more susceptible to plastic flow.


resisting corrosion than coarse-grained additions. However, spalling resistanceis improved with coarse (300 mm) spinel addition, since fine additions resultin excessive sintering shrinkage and consequent stress.82

14.7 Thermal shock

A major factor determining the importance of ceramics lies in their usefulnessat high temperatures. Two types of experimental techniques are generallyused in order to determine the minimum shock required to nucleate fracture(cracking), and the amount of the damage caused by thermal shock:83 thenumber of quenching cycles, and strength as a function of quench temperatureand crack patterns.

One of the common techniques for evaluating thermal shock resistance isthe thermal cycling of materials until fracture occurs. The strength of thematerials after thermal cycling is then compared with the original strength.The number of cycles necessary to cause a defined damage or weight loss isused as a measure of thermal shock resistance. The results of a thermalcycling test can be useful only when used with another thermal shock test,which indicates the relative degree of difficulty of crack initiation andpropagation.84

The traditional standardised method used to characterise the thermal shockresistance of commercial refractory materials normally deals with quenchingof bricks from high temperatures by using water, oil or air as a coolingmedium. A slow decline in strength occurs in weaker, low Young’s modulus,materials; however, a sharp drop in strength is shown by a strong materialsubjected to greater than the critical shock.85 The critical quench temperature,DTc, can be defined as the temperature drop required to produce cracking inhalf the specimens tested. Cracking is usually detected by the dye-penetrantmethod. The depth of cracking increases slowly with increasing quenchingtemperature, and this gives rise to a gradual fall-off in strength. The deepestcracks control the strength after quenching.83

14.7.1 Thermal shock parameters

On the basis of these tests, two types of parameter have been proposed:thermal stress resistance, and thermal shock damage.52,86 Since the term‘thermal shock resistance’ is used to describe both the nucleation of fractureby the thermal shock and the degree of damage by thermal shock, it isproposed here to refer to the resistance to nucleation by thermal shock as‘thermal shock fracture resistance’ or ‘thermal stress resistance’ and to theresistance damage by thermal shock as ‘thermal shock damage resistance’.52,86

The first method expresses the difficulty of crack initiation. The secondexpresses the degree of possibility for further damage by crack propagation


after crack nucleation is measured. Correlations were then sought betweenthese parameters and the observed thermal shock behaviour. It is clear thatthese methods are useful only for evaluating the relative degree of damage(i.e. the fraction of retained strength) by thermal shock.84

Thermal stress resistance parameters

For material initially undamaged, the appropriate parameter expressing thetendency for cracks to be developed, and therefore strength to be lost, can beconsidered to be that for crack initiation. This has been expressed in terms ofthermal stress resistance parameters.25,30,52,86–88 Kingery used the infiniteslab symmetrically heated or cooled with a constant heat transfer coefficientto derive thermal shock fracture resistance parameters R, R¢ and R¢¢ using theequations:

RE

= (1 – )fs na (14.1)

¢Rk

E =

(1 – )fs na (14.2)

¢¢RA

E =

(1 – )fs na (14.3)

where sf is the strength (normally taken to be the bend strength), E is theYoung’s modulus, a is the mean thermal expansion coefficient of the composite,n is Poisson’s ratio, k is the thermal conductivity, and A is a stress reductionterm.

The parameter R is applicable for the case of instantaneous change insurface temperature (infinite h) for conditions of rapid heat transfer; R¢ is fora relatively low Biot modulus (b < 2) for conditions of slow heat transfer; R¢¢is for a constant heating or cooling rate.88 R defines the minimum temperaturedifference to produce fracture under conditions of infinite heat-transfercoefficient, i.e. A = 1. The parameter R is inversely proportional to a. A lowvalue of a is therefore essential for good thermal stress resistance. Thecoefficient of thermal expansion normally increases with increasingtemperature; however, thermal conductivity decreases.

Hasselman’s approach89 to thermal shock fracture is based on the conversionof released elastic energy into surface energy, which gives the thermal conditionsfor fracture initiation. The length of the crack and the conditions are alsoimportant for further crack propagation. Resistance to crack initiation can bemaximised by achieving high strength and thermal conductivity, with lowvalues of thermal expansivity and Young’s modulus. However, avoiding thermalfracture by increasing strength in order to make initiation difficult is dangerousbecause once initiated, crack propagation will be fast and catastrophic.


Thermal shock damage resistance parameters

Hasselman approached the problem of thermal shock resistance to damageby comparing the degree of damage in materials fractured by thermal shock,rather than fracture initiation. Once cracking is initiated, the maximum surfacearea (Smax) of the fracture face is limited by Smax £ U/gWOF, where U is theelastic stored energy per unit volume and gWOF is the effective surface energyor work of fracture per unit projected area of fracture face. He suggested thata thermal shock damage or toughness parameter involving U/gWOF should beuseful for comparing amounts of cracking.83 The Griffith view is that a crackwill start propagating and will continue to propagate as long as the elasticenergy released from the stress field surrounding the crack is greater than thefracture surface energy.90 The extent of crack propagation is directlyproportional to the elastic energy stored at fracture. In contrast, the R¢¢¢ andR≤≤ parameters (see below) are inversely proportional to the elastic strainenergy, and directly proportional to the effective surface energy.88

Hasselman derived the thermal shock damage resistance parameters R¢¢¢and R¢¢¢¢, expressing the ability of the material to resist crack propagationand further damage and loss of strength on thermal shocking:52,86

¢¢¢ ◊R E = 11 –

f2s n (14.4)

R¢¢¢¢ = 1 –

f2

WOFEs

gn◊ (14.5)

The parameter R¢¢¢¢ can be applied to compare the degree of damage ofmaterials with widely different values of gWOF, such as brittle and ductilematerials. R¢¢¢ can also be used to compare the relative degree of damage ofmaterials with similar crack propagation properties, i.e. the same values ofgWOF.91 R¢¢¢ is a simplified formula derived from R¢¢¢¢ by eliminating the termof gWOF energy.

The criteria for minimising the extent of crack propagation, and for obtaininga low degree of damage,84 are high values of the Young’s modulus, Poisson’sratio and surface energy, and low values of strength. For example, when abody with near-zero strength undergoes only a minimal amount of damage,it will still have a low value of strength after thermal shock, simply becauseof its initial low value of strength before thermal shock. Although highvalues of the thermal shock damage resistance parameter R¢¢¢ and R¢¢¢¢ aredesirable, it is clear that these parameters cannot be maximised by letting thestrength (sf) approach zero. There must be some intermediate value of strengthand a resulting degree of damage such that the strength (after thermal shock)remains within acceptable limits.


14.8 Mechanical properties and thermal shock

behaviour of magnesia–spinel composite

refractory materials

There have been a number of reports regarding the similarities and differencesof the fracture toughness of single-crystal (KS) and polycrystalline (KP)magnesium aluminate spinel ceramics. Polycrystalline values are larger thansingle-crystal values, e.g. KP may be 2–10 times KS.92 The values of single-crystal bars gave overall averages of 1.0 ± 0.2 MPa m1/2. Toughness ofpolycrystalline MgAl2O4 was measured by various techniques; indentationfracture and various notched beam values range from 1.4 to 2.2 MPa m1/2,averaging ~1.9 ± 0.2 MPa m1/2.93 The polycrystalline to single-crystal fracturetoughness (KP/KS) ratios are in the range > 1.5 to < 2.5 for Al2O3 and ≥ 2.4to £ 3.2 for MgO.93

The single crystals have a KIC region at approximately 1 MPa m1/2, up to800∞C. However, at high temperatures (>800∞C) they have a second regionof rapidly increasing KIC values (~3 MPa m1/2) with temperature, because ofincreased crack tip dislocation mobility.77 At intermediate temperatures thepolycrystalline data become indistinguishable from single-crystal values;however, above 1000∞C they fall significantly (55%) below the rapidly risingsingle-crystal curves. This suggests in the fracture process zone fracture mayoccur by other mechanisms, such as grain boundary plasticity, which maycause homogeneous crack tip plasticity in the spinel phase.

It is reported1 that fracture toughness and hardness values of sinteredmagnesium aluminate spinel are in the ranges of 1.9–2.7 MPa m1/2 and 9.3–14.9 GPa respectively, with the additions of various amounts of aluminafrom 49% to 75%. For example, in magnesia–spinel refractory materialscontaining 72 wt% Al2O3, the room-temperature fracture toughness valuehas been reported to be 1.9 MPa m1/2. The fracture toughness of dense MgO(r = 3.5 Mg m–3) measured using the single-edge notched beam (SENB)technique is given as ~2.05 ± 0.05 MPa m1/2 for materials with grain sizesranging from 6 to 8 mm.94 Furthermore, thermal conductivity values ofcommercial MgO and spinel are given as ~50 W m–1 K–1 and 15 W m–1

K–1 at 25∞C and ~7 W m–1 K–1 and 5 W m–1 K–1 at 1000∞C, respectively.1 Inaddition, room-temperature values of strength and Young’s modulus forcommercial sintered spinel are also reported as 135 MPa and 260 GPa,respectively.1 The room-temperature strength of the hot-pressed pure MgO95,96

is found to be ~225 and 230 MPa for ~32 and 25 mm MgO grain sizes,respectively. The Young’s modulus of the pure hot-pressed MgO at roomtemperature is stated as ~290 GPa95 (~25 mm grain size), and 258 GPa94

(~7 mm grain size). The room-temperature fracture surface energy of MgOobtained by using SENB (for a grain size ~100 mm) has been reported to be14 J m–2,97,98 and 15 J m–2.99 The fracture surface energy of dense spinel has


been reported to be 4 ± 0.5 J m–2,77 and 5–7 J m–2.78,100 The measured workof fracture (obtained from the areas under the stress-strain curve) is alsoreported to be 35 J m–2 for polycrystalline fully dense MgO (3.56 Mg m–3)of ~100 mm average grain size.101

Some of the thermomechanical properties of MgO–spinel refractories aregiven in Table 14.3,28 though there is no information about thermal shockparameters, fracture surface energy, work of fracture and strength values.For example, the coefficient of thermal expansion decreases almost linearlywith increasing spinel content. Hot strength increases significantly with theaddition of spinel content up to 40% but further additions of spinel in generaldo not change the hot strength values markedly (Section 14.6.4). Young’smodulus falls to a minimum with increasing spinel content at 30–40% addition;however, Young’s modulus increases significantly with spinel content exceeding40% (e.g. back to the MgO value). As the spinel become the continuousphase, Young’s modulus stayed approximately constant between 50 and 80%.28

Thermal shock damage caused by water quenching from 800∞C decreaseswith increasing spinel content up to 40%.28 Compositions with higher than40% spinel show less improvement due to the increase in initial Young’smodulus, in spite of a decreasing thermal expansion coefficient. As shown inTable 14.3, the best combination of properties is assumed to be achieved bythe addition of 40% spinel.28 Using preformed sintered spinel higher than40% leads to a reduction in refractoriness; however, spinel contents less than10% Al2O3 lead to deterioration in thermal shock resistance.102

Various studies based on the industrial work have been made over the last15 years, with the objective of developing magnesia–spinel materials ofimproved resistance to thermal shock and alkali attack. The work done so farhas been mostly phenomenological, and little quantitative understanding ofthe function of the system variables has been developed. It is therefore stilldifficult to specify optimum compositions with confidence, and materialsdevelopment is based largely on trial and error. Much of the magnesia–spinelrefractory currently produced is used for cement kiln linings, where there aretwo conflicting requirements: for thermal shock resistance the optimum spinelcontent should be fairly high; for reduced calcium oxide attack, involvingreaction with aluminium oxide and the formation of low-melting calciumaluminates, the spinel content must be as low as possible. It is clearly necessaryto keep the spinel content to a minimum, while developing the maximumresistance to thermal shock.4 The optimum spinel content has been claimedto be as high as ~40% by previous researchers.46 Subsequent researchershave used a range of model, high purity, magnesia–spinel composites toexamine in detail the effects of spinel particle size and quantity onthermomechanical behaviour, and suggested an optimum compositecomposition for a maximum resistance to further damage by thermal shockof ~20% of spinel.74,75,103 In spite of the importance of this refractory system,

Magnesia–spinel (M

gAl2 O

4 ) refractory ceramic com

posites377

Table 14.3 Properties of impure MgO–spinel systems28

Properties Values of properties at nominal spinel content %

0 10 20 30 40 50 60 70 80

Bulk density (Mg m–3) 3.43 3.42 3.37 3.33 3.31 3.44 3.24 3.36 3.44Apparent porosity (%) 0.0 1.2 1.7 4.3 3.5 0.0 0.0 0.0 0.0Theoretical density (%) 95.8 95.6 94.3 93.2 92.7 96.5 90.9 94.4 96.7Hot strength (3-point loading at 1600∞C)

Failure stress (MPa) 7.2 4.6 6.7 9.2 11.6 11.3 9.8 11.0 12.8Coefficient of thermal expansion (MK–1)

20–1000∞C 13.9 13.1 12.5 11.7 11.5 n.d* 9.6 9.6 9.120–1500∞C 14.9 14.3 13.3 11.7 11.9 n.d* 10.0 9.7 9.5

Young’s modulus (E) at 20∞C (GPa) 274 264 194 151 152 256 213 239 247Thermal shock resistance

Progressive quench test % drop in E 80.8 72.0 38.3 24.8 22.7 58.9 41.7 44.5 54.6quenching from 800∞C into water

*n.d. = not determined.


few fundamental underpinning scientific studies have been reported so far,apart from the most recent researches.45,71,73,74,103,104 It has already beenreported that the predicted thermal shock parameters103,105 and gWOF/gi ratios105

are compared with the actual thermal shock behaviour, and the R¢¢¢¢ parameterand gWOF/gi ratios are found to be good indicators for a quantitative evaluationof the retained strengths of MgO–spinel composites. It is also stated thatresistance to thermal shock damage in terms of the effect of particle sizedistribution of spinel particles on the basis of thermal shock parameters canbe more strongly favoured with materials containing a significantly broaderdistribution of spinel particles, rather than with narrowly distributed spinelparticles for which a much larger volume percentage is required to achievea similar improvement.106

The precise conditions required for the generation of microcracks in thissystem, their mechanism of operation, relationships between microstructureand those mechanical properties likely to affect thermal shock behaviour,through a detailed examination of hot-pressed model composite materials,will be examined in depth in the following sections of this chapter. Emphasisis placed on establishing the nature and extent of microcrack development,and the relationships between composition and microcracking, and betweenmechanical properties and thermal shock behaviour. In general, attention isgiven to the different methods of forming the spinel (as in situ and preformed),including with the different sizes of the spinel particles. It is expected thatthis investigation will provide a platform for detailed modelling of bothmechanical properties and thermal shock behaviour and high-temperatureproperties of magnesia–spinel composite materials, and develop guidelinesto allow the optimisation of commercial magnesia–spinel refractorycompositions and microstructures.

14.8.1 Strength and MgO grain size

Figure 14.2 shows that 0.5 mm in situ spinel composites prepared from Al2O3

powder demonstrate a ~20% decrease in strength for up to 10% addition, butfurther increases do not significantly affect strength. Composites preparedfrom the preformed spinel powder (3, 11 and 22 mm) show a more markeddecrease in strength with increasing amounts of spinel, for all particle sizes.Although the overall pattern of behaviour is similar to that shown by the insitu spinel composites, there is a significant (~55 to 75%) decrease in strengthof MgO at 30% addition, probably because of more extensive microcrackingin the composites.73,104,105 The larger the spinel particles, the more the declinein strength (Fig. 14.2). Spinel content, but predominantly spinel particlesize, significantly affect the strength.

It is reported45,71 that in all composites the mean grain size of MgOincreases by a factor of 2–3 in the presence of 5–10% spinel; however,


further additions of spinel decrease the MgO grain size to that of pure MgO.The deceleration in grain growth with further additions of spinel (≥ 10%) ismost probably because of the dominance of grain boundary pinning effectsby the spinel grains.45 Both acceleration and deceleration are most markedfor the 0.5 mm spinel particles, which are located at the grain boundaries(Fig. 14.3).

There may be a minor effect of MgO grain size in influencing strength,but crack length, spinel particle size and content seem likely to be the majorfactors influencing strength.33

14.8.2 Young’s modulus

Young’s moduli of spinel composites decrease with additions of spinel, forall particle sizes, where the influence of the spinel is greater, the larger the

0.5 mm 3 mm11 mm 22 mm

0 5 10 15 20 25 30Spinel (%)

250

200

150

100

50

0

Str

eng

th (

MP

a)

14.2 Bend strength as a function of in situ formed (0.5 mm) andpreformed (3, 11, 22 mm) spinel content.

14.3 SEM micrograph: MgO containing 0.5 mm 5% in situ spinel.


particle size (Fig. 14.4). SEM observations show (in Figs 14.3 and 14.5) thatcoarse particles are associated with longer crack formation, compared tofiner particles. The crack lengths also increase with increasing concentrationof spinel.73,104 The greater the extent of cracking, the greater the decrease inYoung’s modulus.

The changes in Young’s modulus can be explained in terms of crackdevelopment and interlinking, and these pre-existing cracks are assumed tobe the result of thermal expansion mismatch between the MgO and MgAl2O4

phases.33 When the volume fraction of spinel increases, more microcrackingoccurs and lower Young’s modulus values are obtained.

0.5 mm 3 mm11 mm 22 mm

0 5 10 15 20 25 30Spinel (%)

300

250

200

150

100

50

0

You

ng

’s m

od

ulu

s (G

Pa)

14.4 Young’s modulus as a function of in situ formed (0.5 mm) andpreformed (3, 11, 22 mm) spinel content.

14.5 SEM micrograph showing the microstructure of a 20% 22 mmpreformed spinel composite.


14.8.3 Fracture toughness and fracture surface energy

Figure 14.6 shows that composites prepared from preformed spinel show ageneral marked decrease in toughness with increasing spinel content for allparticle sizes. Values of KIC of preformed spinel composites decrease withup to 10% addition, and level out at ~1 MPa m1/2 with further additions. Itappears that K1C is certainly much more sensitive to spinel content thanspinel particle size, and within experimental error may be independent ofparticle size. On the contrary, K1C for the 5% 0.5 mm in situ spinel compositesinitially decreases, but further spinel additions bring the values back to thoseof pure MgO (Fig. 14.6). This can be attributed to the special microstructureof this material in which the spinel particles were located only along theMgO grain boundaries (Fig. 14.3), as compared to preformed spinel composites.(Fig. 14.5). The significant change in fracture path from transgranular tointergranular mode (explained in Section 14.8.5) is possibly the explanationfor the marked increase in K1C (e.g. back to the MgO value) at higher spinelcontents.45 As will be expected, this pattern of behaviour in fracture surfaceenergy values calculated using experimental Young’s modulus values in generalis found to be similar to that of K1C.45

14.8.4 Critical defect size

The mean critical defect size (c), calculated from the Griffith equation,107,108

increases markedly with increasing spinel content (Fig. 14.7). For 0.5 mm insitu spinel composites, the defect size increases up to ~20%, and then decreasesslightly with further additions of spinel. However, for the preformed spinelcomposites there is a gradual increase in defect size up to 30% spinel additions.The increase in defect size is in general much more marked at 20%, in

0.5 mm 3 mm11 mm 22 mm

0 5 10 15 20 25 30Spinel (%)

2.5

2.0

1.5

1.0

0.5

0

KIC

(M

Pa.

m1/

2 )

14.6 Fracture toughness (KIC) as a function of in situ formed (0.5 mm)and preformed (3, 11, 22 mm) spinel content.


comparison to the pure MgO, and the defect sizes of the in situ formed 0.5mm spinel and preformed 22 mm spinel composites increase approximatelyby factors of 2 and ≥ 1.5 respectively. The main result obtained from Fig.14.7 is that the general trend is for the defect size to increase with additionsof spinel. The highest increase in the defect size of 0.5 mm composites ispossibly because of a grain boundary pinning effect of the spinel particleslocated at the MgO grain boundaries (Fig. 14.3).

It is suggested45,71 that spinel content (predominantly) and spinel particlesize (weakly) are important factors determining the critical defect size. MgOcleavage, grain boundary strength, pre-existing cracks controlling effectivegrain boundary energy, and pores, may be other parameters determiningcritical flaw size. However, particle interaction coupled with thermal expansionmismatch causing microcracking and interlinking that leads to a markedincrease in the critical defect size, seems most likely to be the importantparameter.71

14.8.5 Work of fracture

For the in situ formed 0.5 mm spinel composites, the work of fracture (gWOF)value increases only slightly with up to 10% additions, but at 20% it increasesby a factor of 2.5 (Fig. 14.8). There is a general but less marked increase ingWOF, by a factor of ~2, at 30% additions of preformed spinel (Fig. 14.8). Itseems that the effect of preformed spinel content is more important thanparticle size, within the larger scatter data.

Fracture surfaces of pure MgO show a large proportion of transgranularcracks, with a few intergranular cracks (Fig. 14.9). At low additions of spinel,transgranular cracks are still present with some intergranular cracks, in thefracture surfaces of each spinel composite. However, at higher additions ofspinel (≥ 20%), mostly intergranular fracture occurs. For example, 20%

0.5 mm3 mm11 mm22 mm

60

50

40

30

20

10

0

c (m

m)

0 5 10 15 20 25 30Spinel (%)

14.7 Critical defect size (c) as a function of in situ formed (0.5 mm)and preformed (3, 11, 22 mm) spinel content.


22 mm spinel composite (Fig. 14.10) shows mostly intergranular fracturewith some transgranular. It therefore appears that higher values of gWOF areassociated with the occurrence of more intergranular fracture with increasingspinel additions that requires more energy to propagate a crack completelythrough the specimen.

The fracture of the magnesia–spinel composites is either semi-stable orstable, but never catastrophic.105 It may be concluded that crack propagationis a much greater energy-consuming process than crack initiation in thesematerials. For many industrial applications, the initiation of fracture is less


0 5 10 15 20 25 30Spinel (%)

100

80

60

40

20

0

Wo

rk o

f fr

actu

re (

J m

–2)

14.8 Work of fracture (gWOF) as a function of in situ formed (0.5 mm)and preformed (3, 11, 22 mm) spinel content.

14.9 Fracture surface of dense MgO.


important than gWOF and the degree of damage (e.g. further loss of mechanicalintegrity by strength and material loss through large-scale fracture behindthe hot face).109 Large values of the gWOF/gi ratio are obtained in thesecomposites.105 This is a basic requirement for refractory materials to showgood thermal shock damage resistance.110

14.8.6 R and R ¢¢¢¢ parameters

Figure 14.11 shows that the R parameter for 22 mm spinel composites decreaseswith additions of up to ~20%, and then increases with further additions of

14.10 Fracture surface of composite containing 20% 22 mmpreformed spinel.

0.5 mm 3 mm11 mm 22 mm

80

60

40

20

0

R p

aram

eter

(K

)

0 5 10 15 20 25 30Spinel (%)

14.11 R parameter as a function of in situ formed (0.5 mm) andpreformed (3, 11, 22 mm) spinel content.


spinel back to the value for pure MgO, within experimental scatter. For thecomposites containing lower particle sizes of spinel, the R parameter initiallydecreases slightly with spinel content, and reaches a minimum at about 10%addition: further additions of spinel result in increases in the R parameter.For example, 30% of 0.5 mm in situ formed spinel composite gives an Rparameter ~50% greater than for pure MgO.

It appears that both strength and Young’s modulus are controlled by theextent of microcracking with strength being influenced more strongly, untilvery high spinel contents are reached. It cannot be expected on this basis thatspinel additions will improve the thermal shock resistance of MgO throughimprovement of resistance to crack initiation.74,103

Refractories are not very resistant to crack initiation, but have a significantresistance to crack propagation or extension. Crack propagation is muchmore difficult in refractories than the initiation of cracks.111 The fracturemechanism in magnesia–spinel refractories relies on the development ofmicrocracks that allows easy crack initiation but makes propagation, in whichfracture occurs in a quasi-static manner, more difficult.46 The main concernby assessing the service performance of materials (e.g. strength loss) is theresistance to crack propagation and to extension of damage caused by thermalshock, rather than resistance to crack initiation.112–115 Therefore, there is abasic requirement to investigate how the R¢¢¢¢ parameter varies for eachcomposition.

Figure 14.12 shows that R¢¢¢¢ increases with additions of 0.5 mm in situformed spinel to a maximum at 20% loading, and an improvement by afactor of ~4, as compared to MgO. The R¢¢¢¢ parameter for the 3 mm and11 mm preformed spinel composites shows similar values to the 0.5 mmspinel composites up to 20% additions, but is more sensitive to particle sizewith further additions. The 22 mm preformed spinel composites shows asimilar trend but there is a larger effect on R¢¢¢¢ above 10% additions, as


2.5

2.0

1.5

1.0

0.5

0

R¢¢¢¢ p

aram

eter

(m

m)

0 5 10 15 20 25 30Spinel (%)

14.12 R¢¢¢¢ parameter as a function of in situ formed (0.5 mm) andpreformed (3, 11, 22 mm) spinel content.


compared to the other composites. It will be expected that the 20% 22 mmspinel composite will have the highest resistance to thermal shock damage,with an improvement by a factor of ~9, as compared to pure MgO. Thisincrease in R¢¢¢¢ parameter appears mainly to be the result of the markedincrease in gWOF with spinel additions (Fig. 14.8), and the greater sensitivityof strength to composition and particle size than Young’s modulus.103 Basicallythe R¢¢¢ and R¢¢¢¢ parameters exhibit similar patterns in terms of variations inspinel content.33,74

On the basis of the R plot (Fig. 14.11) spinel composites in general shouldnot be more resistant to crack initiation than pure MgO. As can be expectedon the basis of the R¢¢¢ and gWOF values, the R¢¢¢¢ values predicts that compositescontaining the coarser spinel particles should in general be best at resistingfurther thermal shock damage.33 The fracture mechanism in magnesia–spinelcomposites therefore appears to rely on the development of microcracks;though the composites are not more resistant to crack initiation than is MgO,they have a stronger resistance to crack extension.74

The thermal shock damage resistance of 0.5 mm in situ formed compositeswill be expected to be much less than that of preformed composites. For 3mm and 11 mm spinel composites, addition of 30–40% spinel may providesome improvements. Further deterioration of strength in the coarsest spinelcomposite as a result of thermal shock should be a minimum at 20% (Fig.14.12). In general coarser (22 mm) preformed spinel powders appear to bemore beneficial than finer powders, but there is no obvious advantage withadditions of more than 30%.74,103

14.8.7 Relative strength

One of the common procedures for evaluating thermal shock resistance is tocompare relative strength values, which are the strengths of bars after quenchingrelative to initial strengths, as a function of quench temperature. This providesa direct indication of resistance to further damage caused by thermal shock.84

Relative strengths of shocked MgO, and preformed spinel composites, areshown in Fig. 14.13, together with the initial strength values of each material.

Figure 14.13 illustrates that values for pure MgO were almost constant upto ~575∞C, but further increases in the quench temperature result in a sharpdecrease (~80%) in strength. Above this quench temperature, the strengthvalues remain almost at the same level until 1000∞C. In contrast, the spinelcomposites for all spinel volume fractions have a higher relative strengththan pure MgO from ~600∞C up to the maximum quench temperature used.The absence of any corresponding change in strength for the 20% 22 mmspinel composite suggests that the maximum defect size is unchanged. The20% 22 mm spinel composite does not lose its strength further after thequench tests, and in this respect is the most stable of the composites tested.


Spinel composites should be more useful than pure MgO in terms ofresistance to thermal shock damage and further loss of strength, because therelative strengths of all composite materials after shocking remain muchhigher than those of MgO (Fig. 14.13). These results show clearly that theresistances to thermal shock damage for all these materials have the trendsexpected on the basis of the calculated R¢¢¢ and R¢¢¢¢ parameters.74,103

14.8.8 Degree of damage in MgO and spinel composites

The amounts and distributions of cracking after quenching in fully denseMgO and in composites containing 20% 22 mm spinel are shown in Figs14.14(a) and 14.14(b), respectively.

Observations of pure MgO quenched into oil from <600∞C, on polishedsurfaces and cross-sections through the centres of the bars, show no cracks;however, at the critical quench temperature (DTc), ~600∞C, a large number ofcracks appear on the polished surfaces, and a small number progress towardsthe centres of the bars (Fig. 14.14(a)). The amount of cracks increases markedlyat quenching from 1000∞C, and cracks are now visible on both polishedsurfaces and cross-sections through to the centres of the bars; most crackspenetrate more than one-quarter of the distance through the bar.

Figure 14.14(b) demonstrates that at a quench temperature of 600∞C for20% 22 mm spinel composite, cracks start developing, and the number onpolished surfaces increases smoothly up to a quench temperature of 1000∞C.Cracks do not progress into the bar centres, and the distance of crack penetrationthrough the bar remains very small, even at the highest quench temperatureof 1000∞C.

No particular quench temperature at which crack density increased markedlyis observed using dye penetrant for the spinel composites, in contrast to theMgO materials. It is clearly seen that there is a gradual development of adense microcrack network in the spinel composites, with increase in thequench temperature.

MgO – 233 MPa20% 22 mm – 65.1 MPa20% 11 mm – 111 MPa

0 200 400 600 800 1000Quench temperature (∞C)

1.2

1.0

0.8

0.6

0.4

0.2

0

Rel

ativ

e st

ren

gth

14.13 Relative strengths of pure MgO and preformed spinelcomposites, as a function of quench temperature.


14.8.9 Comparison of fraction retained strength and R ¢¢¢¢parameter

In order to compare the thermal shock damage resistance of MgO and MgO–spinel composites, all the values of strength after shock are compared toinitial strengths, i.e. retained strength after shock, rather than absolute strengthvalues.74 The calculated R ¢¢¢¢ parameter for the MgO and spinel compositesshows a good correlation with retained strength after thermal shock testing(Fig. 14.15).

Figure 14.15 for a quench temperature of 800∞C shows that for pure MgOthere is very severe mechanical property degradation at low values of R¢¢¢¢:

14.15 Fraction retained strengths of MgO, and spinel composites,after a quench from 800∞C, as a function of R¢¢¢¢ parameter.

0 1 2 3R ¢¢¢¢ (mm)

MgO20% 11 mm Spinel10% 22 mm Spinel20% 22 mm Spinel30% 22 mm Spinel

1.0

0.8

0.6

0.4

0.2

0Frac

tio

n r

etai

ned

str

eng

th

2 mm

600∞C

1000∞C

(a)

2 mm(b)

600∞C

1000∞C

14.14 (a) Thermal shock crack patterns revealed by dye penetrant, onthe polished surfaces and cross-sections of MgO bars. (b) Thermalshock crack patterns revealed by dye penetrant, on the polishedsurfaces and cross-sections of 20% 22 mm preformed spinel bars.


pure MgO retained only ~20% of its strength. For the best (20–30% 22 mm)magnesia–spinel composites, the R¢¢¢¢ parameter is approximately five timesthat of the pure MgO. The other composites retain about half of their originalstrength. There is a clear and approximately linear decrease in the extent ofthermal shock damage, with increasing R¢¢¢¢. The retained strength reaches amaximum and levels out with increasing R¢¢¢¢ parameter above the value of~2 mm.

This comparison shows the general relationship that materials with largervalues of the damage resistance parameter, R¢¢¢¢, suffer less thermal shockdamage, and have higher retained strength after quenching. This investigationshows that the R¢¢¢¢ parameter, calculated from strength and Young’s modulusvalues measured at room temperature, can be used to predict the loss instrength of MgO and spinel composites, after thermal shock.74 The observedchanges in strength with thermal shock for these materials are in accord withthose predicted.

14.9 Conclusions

The coarser spinel particles cause greater loss of strength: spinel content,and more importantly spinel particle size, are factors determining strength.The Young’s modulus values of spinel composites in general decrease withincreasing spinel content. Young’s modulus appears to be more sensitive tospinel content than to spinel particle size.

The relationships between the KIC and fracture surface energy are similarto each other, and generally decrease with increasing spinel content, apartfrom the in situ formed composites which show different microstructuresand fracture paths. It is therefore clear that the development of thermal shockresistance in the magnesia–spinel composites cannot be linked to any increasedfracture toughness.

The calculated critical defect sizes, and also actual crack lengths, increasewith spinel additions. MgO grain size varies with spinel content but thisappears to have little if any influence on strength. The thermal expansionmismatch, causing microcracking, longer interlinked cracks, and spinel particlesize and content, are likely to be the major factors determining strength.

There is a general increase in gWOF with increasing spinel additions, andthe effect of spinel content is greater than that of particle size. The R parametergenerally decreases with spinel content for all particle sizes. Magnesia–spinel composites are not more resistant to crack initiation than MgO, becauseof the pre-existing cracks, but they show greater resistance to further crackpropagation than MgO. For this reason, predictions made from the calculatedR¢¢¢¢ parameter must be taken into account in characterising thermal shockdamage and retained strength of materials, rather than the crack initiationresistance parameter, R.


In the MgO – spinel composites a large amount of cracking is createdduring cooling, as a result of high tensile hoop stresses generated around thespinel particles, due to the large thermal expansion difference between MgOand spinel. The improved resistance to thermal shock observed in thesemodel magnesia–spinel composites can therefore be attributed to the microcracknetworks developed around the spinel particles.

Finer spinel particles result in shorter initial crack propagation distancesfrom the spinel particles; the coarser spinel particles are the origins of longercracks. It might therefore have been expected that resistance to thermalshock damage, resulting from resistance to microcrack propagation andinterlinking, would be higher with materials containing the coarser particlespinel.

This prediction can be confirmed by the R¢¢¢¢ parameter for these modelcomposite materials, which suggests an optimum composition of 20% 22 mmspinel particles, indicating that resistance to further thermal shock damage isabout nine times higher than in MgO. The predictions are similar to theresults obtained from strength loss values. The relationships between the R¢¢¢and R¢¢¢¢ parameters are similar to each other, and one is as good as the other;they are good indicators for a quantitative evaluation of the retained strengths.

After quenching at ~600∞C, a large number of cracks appear on the polishedsurfaces of MgO bars, and a small proportion of cracks progress towards thecentres of the bars. In contrast, the 22 mm spinel composites do not show acritical quench temperature: a gradual development of microcrack networksis observed with increasing quench temperature, and cracks do not progressinto the cross-section of the bars.

For pure MgO, cracks propagate catastrophically after shocking fromabove the critical quench temperature, with a large decrease in strength.However, spinel composites of all types have higher relative strengths afterquenching than pure MgO.

The 20% 22 mm spinel composites show a significant improvement inshock damage in terms of increased difficulty of crack propagation. In general,coarser spinel particles are more beneficial than fine, but there is no advantagewith additions of >30% for all the particle sizes.

14.10 Future trends

The multiple disadvantages of bricks containing high (>90%) levels ofmagnesia, and the intrinsic hydration susceptibility of dolomite-based products,seem likely to limit their future application to the specific kiln areas wherethe conditions pertaining do not expose their limitations.11 Spinel is added tomagnesia so that the brick is able to cope with the mechanical stressesexerted during kiln operation. Stress factors prevailing in cement kilns, suchas thermal shocks, thermal expansion and kiln shell ovality, demand high


elasticity from bricks – behaviour that a pure magnesia brick is not able tooffer.4 Magnesia–spinel bricks are more resistant to reducing atmosphere,attack by sulphur oxides, CO2, and alkali compounds than are magnesia–chromite and dolomite bricks, which lack resistance to moisture, leading todestruction of the brick due to an increase in volume.4 With these properties,magnesia–spinel brick is the most suitable brick for peripherical areas of thesintering zone, which lacks a protective coating. Because of its better thermalshock resistance combined with good abrasion strength, it can also berecommended for the discharge zone to nose ring kiln sections.63

The magnesite–chrome and chrome–magnesite ranges of refractory materialsare also intensively used for applications requiring a high hot-strength andresistance to attack by basic slags and molten metals.10 An important exampleof this type of environment is the rotary cement kiln lining, with centre zonetemperatures exceeding 1600∞C and the presence of semi-liquid and corrosivecalcium aluminosilicates. It is well known that increasing concern over thetoxicity of the Cr(VI) produced from Cr2O3 under alkaline conditions hasmeant that the handling and disposal of waste refractories containing chromeare now subject to strict European Union regulations,117 and alternativerefractory materials which do not contain Cr2O3 are needed. Magnesite anddolomite refractories are satisfactory in many respects for the types ofapplication that previously used magnesite–chrome materials, but lack theirgood resistance to thermal shock. It has been found that this deficiency canbe overcome by incorporating into a magnesite matrix 9–30% of particulatemagnesium aluminate spinel (MgAl2O4), to form the magnesia–spinelmaterials.12,28 Magnesia–spinel refractories were first evaluated more than30 years ago,10 but it has only been during the last decade that marked effortshave been made to use them as alternative refractories to magnesia–chromematerials. Fortuitously, magnesia–spinel refractories also have the additionalimportant advantage of a longer (1.5–2 times) life than an otherwise equivalentmagnesite–chrome refractory, particularly in locations where they are exposedto high temperatures and severe thermal shock.53

Various studies based on the industrial work have been made over the last15 years, with the objective of developing magnesia–spinel materials withimproved resistance to thermal shock and alkali attack. The work done so farhas been mostly phenomenological, and little quantitative understanding ofthe function of the system variables has been developed. It is therefore stilldifficult to specify optimum compositions with confidence, and materialsdevelopment is based largely on trial and error. Much of the magnesia–spinelrefractory currently produced is used for cement kiln linings, where there aretwo conflicting requirements: for thermal shock resistance the optimum spinelcontent should be fairly high; for reduced calcium oxide attack, involvingreaction with aluminium oxide and the formation of low-melting calciumaluminates, the spinel content must be as low as possible. It is clearly necessary


to keep the spinel content to a minimum, while developing the maximumresistance to thermal shock.4 The other technical problems still to be solvedin the magnesia–spinel system are the two most difficult problems of achievingboth good clinker adhesion and low thermal conductivity.63 These must beachieved without sacrificing resistance to chemical and thermo-physical damageand at a cost comparable with other conventional materials. The processparameters such as chemistry, granulometry and design of products producingchanges in porosity, texture, strength and heat-transfer characteristics mustbe investigated in order to obtain maximum cost-effectiveness.63 The systemmost capable of long-term development combines magnesia with preformedspinel.

It has also been suggested118,119 that mechanical properties and thermalshock of magnesia–spinel refractory materials may be improved by addingsmall amounts of TiO2 or ZrO2, but full details have not been given. Thisconclusion needs to be confirmed by further work, and any basic underlyingmechanism established. Thermal shock and mechanical tests should be carriedout on fully characterised materials to determine the optimum amount ofadditives and their composition, and the optimum MgO grain size, for thebest thermal shock resistance.


The contributions of Dr P.D. Warren, who provided suggestions and guidanceregarding the investigation of the study, and Professor B. Rand, are gratefullyacknowledged. P. Bartha, S. Plint, and M.W. Roberts are thanked for helpfuldiscussions. The contributions of the late Professor R.W. Davidge to theplanning of this investigation are acknowledged. The authors also wish toacknowledge P. Bartha,2,4 H.J. Klischat,4 S.C. Cooper28 and P.T.A. Hodson28

for all published sources used by modifying in this chapter (Fig. 14.1 andTables 14.1, 14.2 and 14.3).

14.12 Sources of further information

∑ Alcoa Corporate Center, 201 Isabella Street, Pittsburgh, PA 15212-5858,USA http://www.alcoa.com

∑ Baker Refractories, Steetley Works, Nottinghamshire S80 3EA, UK http://www.emnet.co.uk/baker-refractories/

∑ Calbex Mineral Trading Inc., Wenhua 5#, Zhengzhou, Henan 450000,China http://www.calbex.com

∑ Capital Refractories Ltd, Station Road, Clowne, Derbyshire S43 4AB,UK http://www.capital-refractories.com

∑ C.E. Minerals, 901 East 8th Avenue, King of Prussia, PA 19406, USAhttp://www.ceminerals.com


∑ Coating and Crystal Technology, RD #4 Box 113-B, Cadogan Road,Kittanning, PA 16201, USA http://www.coatingandcrystal.com

∑ Custom Technical Ceramics, Inc., 8041 North I-70 Frontage, Unit 6 Arvada,CO 80002, USA http://www.customtechceramics.com

∑ China Industrial Resources Co Ltd, No. 64-302, Wanlian Villa, 2nd Ave.TEDA, Tianjin 300457, China http://www.circogroup.com

∑ Didier Werke AG, RHI AG, Wienerbergstraße 11, A-1100 Vienna, Austriahttp://www.rhi-ag.com

∑ Magneco/Metrel, Inc., 223 Interstate Road, Addison IL 60101, USA http://www.magneco-metrel.com

∑ Refractarios Peruanos S.A., Casilla 2828, Lima, Peru http://www.refractoriesinstitute.org

∑ Refratechnik Cement GmbH, Rudolf Winkel Strasse 1, D 37079, Göttingen,Germany http://www.refratechnik.com/

∑ Resco Products, Inc., 2 Penn Center Blvd. Suite 430, Pittsburgh, PA 15276,USA www.rescoproducts.com

∑ Roskill Information Services, 27a Leopold Road, London SW19 7BB,UK http://www.roskill.com/

∑ Saint-Gobain C.R.E.E., Research and Development Center, 550 AvenueAlphonse Jauffret, BP 224, 84306 Cavaillon, France http://www.refractories.saint-gobain.com

∑ Sanac SPA, Viale Certosa 249, 20151 Milano, Italy www.sanac.com∑ Tiger Industrial Ceramics Co. Ltd, Zhangli Village, Xiangdong Town,

Pingxiang City, Jiangxi Province, China http://www.alibaba.com∑ Vesuvius Dyko GmbH, Group Sachon, Wiesenstrasse 61, 40549 Düsseldorf,

Germany http://www.sachon-exportadressbuch.de∑ Vesuvius Group sa/nv, Mechelsesteenweg 455/1, 1950 Kraainem, Belgium

http://www.vesuvius.com/index.htm∑ Vrag (Veitsch Radex AG), RHI Refractories, Technology Center, RHI

AG, Magnesitstr. 2, A-8700 Leoben, Austria http://www.veitsch-radex.com∑ VRW Refractories, A Division of South India Corporation (Agencies)

Ltd, 1513 GIDC, Kerala Industrial Estate, near Bavla, Ahmedabad -382 220, 600 095, Tamil Nadu, India http://www.vrwrefractories.com

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59. Bilham, M.A., ‘Magnesia–spinel: the chrome free solution’, International CementReview, 1991 40–1.

60. Uchikawa, H., Hagiwara, H., Shirasaka, M. and Watanabe, T., ‘Application of periclase–spinel bricks to cement rotary kiln in Japan’, Interceram. Spec. Issue, 1984 33 386–406.

61. Macey, C.L., ‘Evaluation of magnesite–spinel refractories for mineral processingkilns’, Industrial Heating, 1992 28–9.

62. Olbrich, M., ‘Fully automated thermal shock test method for testing fired refractorybrick’, Radex-Rundschau, 1990 (2/3) 268–74.

63. Hobrecht, E.J., Daldrup, H.G. and Bartha, P., ‘Development in basic bricks’, Cements,Betons, Platres, Chaux, 1988 (773) 219–25.

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69. Kingery, W.D., ‘Factors affecting thermal stress resistance of ceramic materials’, J.Am. Ceram. Soc., 1955 38(1) 3–15.

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77. Stewart, R.L. and Bradt, R.C., ‘Fracture of polycrystalline MgAl2O4’, J. Am. Ceram.Soc., 1980 63(11) 619–23.

78. Stewart, R.L. and Bradt, R.C., ‘Fracture of single crystal MgAl2O4’, J. Mater. Sci.,1980 15 67–72.

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81. Cortes, G., ‘Magnesia alumina spinels for the refractory industry’, Ceramic TechnologyInternational, 1994 109–12.

82. Korgul, P., Wilson, D.R. and Lee, W.E., ‘Microstructural analysis of corroded alumina–spinel castable refractories’, J. Eur. Ceram. Soc., 1997 17 77–84.

83. Davidge, R.W. and Tappin, G., ‘Thermal shock and fracture in ceramics’, J. Brit.Ceram. Soc., 1967 66 405–22.

84. Hasselman, D.P.H., ‘Elastic energy at fracture and surface energy as design criteriafor thermal shock’, J. Am. Ceram. Soc., 1963 46(11) 535–40.

85. Morrell, R., Handbook of Properties of Technical and Engineering Ceramics, Part1, Her Majesty’s Stationery Office, London, 1985.

86. Hasselman, D.P.H., ‘Unified theory of thermal shock fracture initiation and crackpropagation in brittle ceramics’, J. Am. Ceram. Soc., 1969 52(11) 600–4.

87. Hulbert, S.F., ‘Models for solid state reactions in powdered compacts’, J. Brit.Ceram Soc., 1968 1–32.

88. Chaklader, A.C.D. and Bradley, F., ‘Thermal shock resistance parameters and theirapplication to refractories’, UNITECR ’89, 1989.

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96. Rice, R.W., ‘Strength and fracture of hot-pressed MgO’, Proc. Brit. Ceram. Soc.,1972(20) 329–63.

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98. Davidge, R.W., ‘The texture of special ceramics with particular reference tomechanical properties’, Proc. Brit. Ceram. Soc., 1972(20) 364–78.

99. Unchno, J.J., Bradt, R.C. and Hasselman, D.P.H., ‘Fracture surface energies ofmagnesite refractories’, Am. Ceram. Soc. Bull., 1976 55(7) 665–8.

100. Swanson, G.D., ‘Fracture energies of ceramics’, J. Am. Ceram. Soc., 1972 55(1)48–9.

101. Clarke, F.J.P., Tattersall, H.G. and Tappin, G., ‘Toughness of ceramics and theirwork of fracture’, Proc. Brit. Ceram. Soc., 1996(6) 163–72.

102. Olbrich, M. and Dobrowsky, F., ‘Periclase spinel bricks in the cement industry’,Zement-Kalk-Gibs, 1989 12 307–9.

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109. Hasselman, D.P.H., ‘Strength behaviour of polycrystalline alumina subjected tothermal shock’, J. Am. Ceram. Soc., 1970 53(9) 490–5.

110. Nakayama, J., Abe, H. and Bradt, R.C., ‘Crack stability in the work-of-fracturetest: refractories applications’, J. Am. Ceram. Soc., 1991 64(11) 671–5.

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114. Coppola, J.A. and Bradt, R.C., ‘Thermal shock damage in SiC’, J. Am. Ceram.Soc., 1973 56(4) 214–17.

115. Ainsworth, J.H. and Herron, R.H., ‘Thermal shock damage resistance of refractories’,J. Am. Ceram. Soc. Bull., 1974 53(7) 533–8.

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117. Bray, D.J., ‘Toxicity of chromium compounds formed in refractories’, Bull. Amer.Ceram. Soc., 1985 64(7) 1012–16.

118. Mosser, J., Buchebner, G. and Dosinger, K., ‘New high-quality MgO–Cr2O3 bricksand Cr-free alternatives for the lining of RH/DH-vessels’, Veitsch-Radex Rundschau,1997 1 11–23.

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400

15.1 Introduction

The aim of this chapter is to present an overview of the performance ofceramic matrix composites (CMCs) under conditions of thermal shock, i.e.when they are subjected to sudden changes in temperature during eitherheating or cooling.

Such conditions are possible in the high-temperature applications for whichthese materials are targeted (e.g. Ohnabe et al., 1999). For example, whilethermal shock is not a concern during steady-state operation of a gas turbine,it becomes of great importance during emergency shut-downs, when cool airdrawn from the still spinning compressor is driven through the hot sectionsand can result in a temperature decrease of more than 800∞C within onesecond at the turbine inlet (Baste, 1993). The fact that such a situation mayarise about 100 times during the lifetime of a modern gas turbine engineshows how important it is to assess, and possibly model, the effect of thermalshock on the mechanical and thermal properties of CMCs. Another examplecomes from the nuclear industry, where SiC reinforced with SiC fibres hasbeen proposed as structural material for the first wall and blanket in severalconceptual design studies of future fusion power reactors (Jones et al., 2002).In this case, apart from the moderate shocks inflicted during start-up andshut-down of the system, the plasma-facing material can suffer rapid heatingdue to plasma discharges.

The description of the thermal shock behaviour of CMCs is given withreference to the thermal shock resistance of monolithic ceramic materials.Monolithic ceramics have greater thermal shock sensitivity than metals andcan even suffer catastrophic failure due to thermal shock because of anunfavourable ratio of stiffness and thermal expansion to strength and thermaldiffusivity, and their limited plastic deformation.

The structure of the chapter is as follows: the stress field developed in athermally shocked component is described in Section 15.2 and maximumstresses are identified and quantified. Section 15.3 contains an overview of

15Thermal shock of ceramic matrix

composites

C K A S T R I T S E A S, P S M I T H andJ Y E O M A N S, University of Surrey, UK

Thermal shock of ceramic matrix composites 401

the experimental methods used to simulate thermal shock conditions in thelaboratory and the means utilised to assess its impact on ceramics and CMCs.The behaviour of monolithic ceramics is described in Section 15.4 withreference to the main mechanical and thermal properties that affect it. Inaddition, methods to model this behaviour are presented. Section 15.5concentrates on the thermal shock behaviour of particle- and whisker-reinforcedCMCs, while Section 15.6 contains an extensive review of damage modessustained in fibre-reinforced CMCs due to thermal shock, their effect onproperties, the role of the interface and attempts to analyse and model thesituation.

It should be noted that this review concentrates on thermal shock (i.e. asingle thermal cycle) and no attempt is made to incorporate and describe theeffects of cyclic thermal loading (cyclic thermal shock, thermal shock fatigue,etc.) on the behaviour of CMCs. For information regarding cyclic thermalloading of ceramics and CMCs the reader is advised to consult the extensivereview of Case (2002). Additionally, recent studies have shown that laminatedceramic–metal systems (Sherman, 2001) and layered ceramic–structures(Vandeperre et al., 2001) exhibit better resistance to thermal shock comparedwith monolithic materials. However, such systems are also beyond the scopeof this contribution.

15.2 Thermal shock of brittle materials:

the induced stress field

When a body is subjected to a rapid temperature change such that non-lineartemperature gradients appear, stresses arise due to differential expansion ofeach volume element at a different temperature. The temperature at eachpoint changes with time at a rate dependent on the coefficient of surface heattransfer (HTC) between the medium of different temperature and the body,the shape of the body, and its thermal conductivity. High HTCs, largedimensions, and low thermal conductivities result in large temperature gradientsand, thus, large stresses. This leads to the establishment of a dimensionlessparameter, the ‘Biot modulus’, for the description of the heat transfer condition(Kreith, 1986):

Bi l hk

= (15.1)

where l is a characteristic material dimension (e.g. the half-thickness of aplate), h is the HTC between the body and the medium, and k is the thermalconductivity of the body. The larger the value of Bi, the larger is the rate ofheat transfer between a medium of different temperature and the body.

The sudden temperature change (DT) that generates non-linear temperaturegradients in a body and, as a consequence, thermal stresses is termed ‘thermal


shock’. If DT is positive (i.e. the temperature change is downward) thematerial is subjected to a ‘cold’ shock, whereas if DT is negative the materialis subjected to a ‘hot’ shock. The term refers to a single thermal cycle (N =1) in contrast to terms such as ‘thermal cycling’, ‘cyclic thermal shock’, and‘thermal fatigue’ that apply to multiple thermal cycles (N > 1).

For the calculation of the thermal shock-induced stresses, we consider theplate shown in Fig. 15.1 with Young’s modulus E, Poisson’s ratio n, andcoefficient of thermal expansion (CTE) a, initially held at temperature Ti. Ifthe top and bottom surfaces of the plate come into sudden contact with amedium of lower temperature T• they will cool and try to contract. However,the inner part of the plate initially remains at a higher temperature, whichhinders the contraction of the outer surfaces, giving rise to tensile surfacestresses balanced by a distribution of compressive stresses at the interior. Bycontrast, if the surfaces come into contact with a medium of higher temperatureT•, they will try to expand. As the interior will be at a lower temperature, itwill constrain the expansion of the surfaces, thus giving rise to compressivesurface stresses balanced by a distribution of tensile stresses at the interior.

If perfect heat transfer between the surfaces and the medium is assumed(i.e. if Bi Æ •) the surface immediately adopts the new temperature whilethe interior of the plate remains at Ti. Following Munz and Fett (1999), thiscase corresponds to having a plate that can expand freely in the z-directionwith suppressed expansion in the x- and y-directions. In the absence ofdisplacement restrictions, the plate would expand along the x- and y-directionsby thermal strains of:

ex = a (T• – Ti) (15.2)

ey = a (T• – Ti) (15.3)

Since thermal expansion in both directions is completely suppressed, elasticstrains are created that compensate the thermal strains, i.e.

eel,x + eth,x = 0 (15.4)

eel,y + eth,y = 0 (15.5)

From equations (15.2)–(15.5) we have:

x

zy

2H

syTS

sxTS

15.1 Schematic of a plate of thickness 2H subjected to thermal shock.


eel,x = –eth,x = –a (T• – Ti) = a (Ti – T•) = aDT (15.6)

eel,y = –eth,y = –a (T• – Ti) = a (Ti – T•) = aDT (15.7)

The elastic strains cause ‘thermal stresses’ along the x- and y-axes and can bewritten as:

e s nsel,

TS TS

= – xx y

E E(15.8)

es ns

el,

TS TS

= – yy x

E E(15.9)

By substituting (15.6) and (15.7) in (15.8) and (15.9) respectively and solvingfirst for s x

TS and then for s yTS we can obtain the thermal shock-induced

stresses along the x- and y-axes as:

s s anx y

E TTS TS = = 1 –

D(15.10)

Equation (15.10) shows that thermal shock induces a biaxial stress field,whose maximum value depends on the elastic properties of the material andthe imposed temperature differential.

However, if the rate of heat transfer is not infinite the thermal shock-induced stresses will gradually build up and after some time reach a peakvalue that will be a fraction of the value given by equation (15.10). Thesolution requires transient stress analysis such as those of Cheng (1951) andManson (1966) with the assumption of the plate of Fig. 15.1 being infinite.Following Lu and Fleck (1998), the plate is initially held at temperature T0

and at time t = 0 its top and bottom faces (at z = ± H) are suddenly exposedto a convective medium of temperature T•. The surface heat flow is assumedto satisfy

k

Tz

h T T z Hz = ( – ), at = ∂∂

±•m (15.11)

where kz is the thermal conductivity in the z-direction and T(z, t) is thetemperature of the material. The plate is assumed to be a uniform, linearthermoelastic solid and is analysed under the constraint that it is free toexpand with vanishing axial force

–

TS

–

TS = = 0H

H

xH

H

ydz dzÚ Ús s (15.12)

and vanishing normal stress in the through-thickness direction, i.e. sz = 0.The transient thermal shock-induced stresses, sx(z, t) = sy(z, t), associated

with the temperature distribution T(z, t) are then given by:


s sx yz t z tTS TS( , ) = ( , )

= –( – )1 –

+ (1 – )2

( – )i

–i

E T T EH

T T dzH

Han

an Ú (15.13)

To obtain the temperature distribution T(z, t) the heat flow in the through-thickness direction needs to be considered. This is governed by:

∂∂

∂∂

£2

2 = 1 , | | Tz k

Tt

z Hz

(15.14)

This equation is solved with heat transfer boundary condition (15.11) by astandard separation-of-variables technique, to give:

T TT T

k tH

z Hz t

n nz n n

n n n

( , ) i

i =12

2

– –

= –1 + 2 exp – sin cos( / )

+ sin cos •

• ÊË

ˆ¯ ¥S b b b

b b b

(15.15)

where bn are the roots of bn tan bn = Bi. The thermal shock-induced stressesare obtained from (15.13) and (15.15), and are written in non-dimensionalform as:

s s sa n

sa nx y

x yz tE T T

z t

E T TTS TS

TS

–1i

TS

–1i

= = ( , )

(1 – ) ( – ) =

( , )

(1 – ) ( – )• •

= –

– – 1

2 –

– ( , ) i

i –

( , ) i

i

T TT T H

T TT T

dzz t

H

Hz t

• •Ú= 2 exp –

sin + sin cos

=1

22S

n nz n

n n n

k tH

• ÊËÁ

ˆ¯

b bb b b

¥ ÊË

ˆ¯

ÏÌÓ

¸˝˛

cos – sin

b bbn

n

n

zH

(15.16)

The evolution of dimensionless stresses is then plotted against dimensionlesstime ( = / )2t k t Hz at selected locations (z/H) through the thickness of theplate and for various values of the Biot modulus. An example of such a plotis given in Fig. 15.2.

The plots show that under cold shock and for all values of Bi, the maximumtensile stress is achieved at the surfaces while the maximum compressivestress is achieved at the centre of the plate. The opposite is true for hot shockconditions. The maximum tensile stress, s max

TS , achieved at the surface duringcold shock and at the centre during hot shock, is then plotted against 1/Bi, asshown in Fig. 15.3.

It can be seen that s maxTS increases with increasing Bi for both cold and hot


15.2 The evolution of dimensionless thermal shock-induced stresswith dimensionless time at various locations through the platethickness for Bi = 10 (reprinted from Acta Materialia, 46, Lu andFleck, ‘The thermal shock resistance of solids’, 4755–4768, copyright1998, with permission from Elsevier).

1

0.8

0.6

0.4

0.2

0

sa

max

1(

–

)E

TT•

NumericalCurve fit

Surface

Centre

0 2 4 6 8 101/Bi

15.3 The maximum tensile thermal shock-induced stress achieved atthe surface in cold shock and in the centre of the plate in hot shockas a function of 1/Bi. Also shown are curve fits expressed describedby equations [15.17] and [15.18] (reprinted from Acta Materialia, 46,Lu and Fleck, ‘The thermal shock resistance of solids’, 4755–4768,copyright 1998, with permission from Elsevier).

Bi = 10

zH

= 1

k t

H2

sa

(,

)

( –

)

izt

ET

T•

0.8

0.6

0.4

0.2

0

– 0.2

– 0.40 0.1 0.2 0.3 0.4 0.5

0.9

0.8

0.7

0.60.5

0.40.2

0


shock, whereas the maximum tensile stress developed at the surface duringcold shock is always much higher than the peak tensile stress developed atthe centre of the plate during hot shock. This observation, combined with thefact that brittle materials usually contain a distribution of surface flaws,means that cold shock is a much more dangerous condition for a brittlematerial.

The maximum surface stress in the infinite plate under cold shock isadequately described by the formula:

s maxTS –16/

–1

( , *) = 1.5 + 3.25 – 0.5e± ÊË

ˆ¯H t

BiBi (15.17)

where t* is the time taken for this value to be reached. The maximum tensilestress at the centre of the plate for hot shock is given by:

s maxTS (0, *) = 0.3085

1 + (2/ )t

Bi(15.18)

Equation (15.17) can be written through (15.10) as:

s anmax

TS –16/–1

= 1 –

1.5 + 3.25 – 0.5E T

Bie BiD Ê

Ëˆ¯ (15.19)

and subsequently as:

s anmax

TS = 1 –

AE TD

(15.20)

In equation (15.20) the parameter A is termed the ‘stress reduction factor’and is given by:

A

Bif BiBi

= 1

1.5 + 3.25 – 0.5e = 1

( )–16/ (15.21)

Equation (15.20) is the classic formula used to characterise thermal shock-induced stresses at the surfaces of brittle components during cold shock. Thefunction f (Bi) can be written more generally as:

f Bi a bBi

c d Bi( ) = + – e / (15.22)

where the values of a, b, c and d depend on the shape of the component andare determined by using analyses similar to the one presented above for aninfinite plate. For example, it was shown that for an infinite plate a = 1.5,b = 3.25, c = 0.5 and d = –16, whereas for an infinite rod a = 1.5, b = 4.67,c = 0.5 and d = –51 (Manson, 1966).

The value of f (Bi) (and consequently A) depends on the Biot modulus, i.e.on the HTC, thermal conductivity and material dimensions. For severe shocks


Bi becomes very large, so f (Bi) becomes f (Bi) ª a, which leads to the maximumvalue of the stress reduction factor being given by A ª 1/a (Wang and Singh,1994).

Experimental evidence suggests that there is a critical value of thecharacteristic specimen dimension, lc, above which Bi (and consequently Aand the value of the shock-induced stress) becomes independent of materialdimensions (Wang and Singh, 1994). Becher and Warwick (1993) showedgraphically that this value may be approximated by:

l ba

khc ª (15.23)

For example, for an infinite rod, the critical dimension is given by lc ª3.1k/h.

15.3 Experimental methods

15.3.1 Introduction

This section aims to present briefly the experimental methods used to evaluatethe performance of ceramics and CMCs under conditions of thermal shock.Reference is made to techniques used to impose the actual thermal shockcondition as well as the destructive and non-destructive methods employedto assess damage morphologies and changes in residual properties.

15.3.2 Thermal shock simulation methods

The methods used to simulate thermal shock can be classified into twocategories, depending on the sign of the temperature differential to which thematerial is exposed:

1. Quench tests, when the material is subjected to a sudden temperaturedecrease (DT < 0)

2. Fast heating methods, if a sudden increase in temperature (DT > 0) isinvolved.

In a quench test, the specimen is heated to a pre-determined temperaturein a furnace and is held at that temperature for a certain period of time (~10–20 min) to allow for the furnace and specimen temperatures to reachequilibrium. A sudden temperature decrease is then brought about by bringingthe heated specimen into contact with a cooling agent. The difference intemperatures between the specimen and the cooling agent is defined as the‘quenching temperature difference’. The process is repeated for differentfurnace temperatures until the temperature at which fracture and/or propertydegradation is just initiated can be determined. The difference in temperature


between this and the medium is the ‘critical quenching temperature difference’,DTc.

Methods to cause rapid temperature decrease of the heated specimeninclude immersing the specimen into a quenching medium, subjecting it toa flow of high velocity cold air (Faber et al., 1981) or water (Absi andGlandus, 2004), and contacting the specimen with a cold metal rod (Rogersand Emery, 1992). The most popular method has been the first, while themost commonly used quenching medium is room-temperature water (Wangand Singh, 1994). Other quenching media include boiling water (Becher,1981; Tiegs and Becher, 1987), room-temperature air (Boccaccini et al.,1998, 1999), glycerine oil (Ishitsuka et al., 1989; Uribe and Baudin, 2003),silicone oil (Evans et al., 1975; Konsztowicz, 1990, 1993; Tancret andOsterstock, 1997), ethylene glycol (Thompson and Rawlings, 1991), methylalcohol (Ishitsuka et al., 1989), liquid nitrogen (Lee et al., 1993; Tancret andOsterstock, 1997), liquid metals (Hencke et al., 1984), pre-heated salt (Soboyejoet al., 2001), or fluidized beds (Morrel, 1993; Schneibel et al., 1998).

The advantages of the quench test include its simplicity and the well-defined temperature difference between sample and cooling agent. However,a major drawback is that the value of the HTC is often difficult to assess,especially for quenching into water where h is affected by different boilingphenomena (Kreith, 1986). In addition, the HTC is not a constant for acertain quenching medium, as it changes with specimen temperature and isaffected by the surface finish of the specimen (Becher et al., 1980; Becher,1981; Becher and Warwick, 1993). Quenching into media other than waterresults in significantly lower values of h (Lee et al., 1993). For these reasons,quenching experiments are suitable for comparing materials but not formeasuring absolute values (Pompe et al., 1993; Morrel, 1993).

In a fast heating test, usually the central area of a specimen is quicklyheated up by a heating source. Heating sources used include plasma jets,lasers, energetic electron beams, hot gas jets, arc discharges and hydrogen–oxygen flames (Pompe et al., 1993; Schneider and Petzow, 1993). The fastheating test causes a different stress distribution in the specimen to thequench test, which results in the activation of a different population of flaws.The thermal shock resistance of a material can be evaluated by measuringthe critical temperature difference for crack initiation (i.e. as in the quenchtest), by measuring the critical power of the heating source for failure, or bymeasuring the temperature gradient as a function of time and calculating thecorresponding energy input and stress intensity factor (Wang and Singh,1994). Generally, the thermal shock induced by a heating source is consideredto be much less severe than that imparted during a quench test, especiallywhen room-temperature water is the quenching medium (Case, 2002).


15.3.3 Assessment of thermal shock damage

The impact of thermal shock on the properties of a ceramic or a CMC isassessed by means of both destructive and non-destructive testing methods.Flexural or tensile (mainly for CMCs) tests of suitably-sized thermally shockedspecimens are usually employed to measure retained mechanical propertiesas a function of the temperature difference. The temperature differential forwhich a significant drop in property values is observed is the DTc. For monolithicceramics and particle- or whisker-reinforced CMCs the property underinvestigation is usually strength, whereas in fibre-reinforced CMCs a drop inYoung’s modulus is usually a better indication of the onset of damage.

Alternative approaches, termed ‘indentation thermal shock tests’, withpre-cracks of known sizes have been used by several authors to assess thermalshock damage in monolithic ceramics. Knoop (Hasselmann et al., 1978;Faber et al., 1981) or Vickers (Gong et al., 1992; Osterstock, 1993; Anderssonand Rowcliffe, 1996; Tancret and Osterstock, 1997; Collin and Rowcliffe,1999, 2000; Lee et al., 2002) indentations were made on rectangular bars,which were then heated to pre-determined temperatures and quenched intowater. Crack extensions from the indentations were measured as a functionof quench temperature differential, and the critical temperature for spontaneouscrack growth (failure) was determined for the material. Fracture mechanicsanalyses, which took into account measured resistance-curve (R-curve)functions, were then used to account for the data trends.

In addition, it has been shown (Boccaccini et al., 2001; Chlup et al., 2001)that the chevron-notched specimen flexural technique (the CN-technique)can be a reliable method of assessing fracture properties (fracture toughness,work of fracture) in thermally shocked brittle matrix composites reinforcedby brittle fibres.

As an alternative to destructive methods, various non-destructive techniqueshave been employed to assess damage caused by thermal shock. These includethe determination of the post-shock Young’s modulus using ultrasonics orthrough the identification of the mechanical resonant frequencies of thematerial (Carter et al., 1988; Lee and Case, 1989, 1990; Wang and Singh,1994; Wang et al., 1994, 1996; Boccaccini et al., 1997, 1998), the monitoringof the change in the spectra of ultrasonic pulses passed through a specimenbefore and after thermal shock (Thompson and Rawlings, 1991), the acousticemission technique (Evans et al., 1975; Konsztowicz, 1990, 1993; Rogersand Emery, 1992), the measurement of the change in specific damping capacity(Q–1) (Lee and Case, 1989, 1990; Boccaccini et al., 1997, 1998, 1999), andthe measurement of thermal diffusivity before and after the shock using the‘flash diffusivity’ method (Ellingson, 1995; Graham et al., 2003).

In addition, optical microscopy (e.g. reflected light microscopy) and scanningelectron microscopy have been used extensively for direct observation of


crack patterns on suitably polished surfaces of thermally shocked samples.More recently, a thermal shock testing technique has been developed(Wereszczak et al., 1999) that uses a high-resolution, high-temperature infraredcamera to capture the surface temperature distribution of a test specimen atfracture.

15.4 Thermal shock of monolithic ceramics

The behaviour of ceramic materials under conditions of thermal shock ischaracterised by a number of parameters (figures-of-merit) that concentrateon either the initiation of cracking due to thermal shock or the resistance ofa material to crack propagation during thermal shock (Kingery 1955;Hasselman, 1970, 1978, 1985). The first parameter is derived by consideringequation (15.10) and assuming that fracture occurs when the thermally inducedstress, sTS, becomes equal to the strength of the material, st. By solving(15.10) for DT (= DTc) we obtain the ‘maximum allowable quenchingtemperature difference’ for the onset of cracking under severe thermal shockconditions, i.e. conditions that approximate perfect heat transfer between thematerial surface and the quenching medium (e.g. in water quench), as:

RE

= (1 – )ts na (15.24)

Equation (15.24) shows that, for high resistance to crack initiation, highstrengths combined with low stiffness and CTE are required. Under mildthermal shock conditions (e.g. in a boiling water quench) the thermalconductivity also becomes important, and (15.24) is modified to give:

¢RE

k = (1 – )ts na (15.25)

The resistance to crack propagation is characterised by the followingparameter:

R≤≤ = (1 – )t

2

EGs n (15.26)

where G is the surface fracture energy. Equation (15.26) shows that forbetter resistance to crack propagation, high values of stiffness and toughnessare required, combined with low strengths.

Different parameters impose different requirements on ceramic materialsdepending on whether fracture resistance or crack propagation resistance isof prime importance. The values of the above parameters for a range ofceramic materials are presented in Table 15.1, where the property dependenceof thermal shock behaviour can be observed. A variety of other parameters


have also been proposed that characterise the thermal shock behaviour ofbrittle materials under a range of different conditions, and these can be foundin reviews such as that of Wang and Singh (1994).

Alternative analyses, aiming to combine the two different approaches,have been performed using fracture mechanics concepts. Hasselman (1969)considered a brittle solid that contained circular, uniformly distributed Griffithmicrocracks. Crack instability due to thermal shock was assumed to takeplace by the simultaneous radial propagation of N cracks of radius l in a unitvolume. Hasselman proposed that the driving force for crack propagation isderived from the elastic energy stored in the body at the instant of fracture.The total energy per unit volume of a body is the sum of the elastic energyand the fracture energy of the cracks, i.e.:

WL T E NL

NLt

2 2 3 –12 =

3( )2(1 – 2 )

1 + 16(1 – )

9(1 – 2 ) + 2

Dn

nn p gÏ

ÌÓ

¸˝˛

(15.27)

Cracks are unstable between those limits for which:

dWdL

t = 0 (15.28)

Combining (15.27) and (15.28), we get the critical quenching temperaturedifference as:

DTG

E LG N L

Ec

2

2 2

1/2 2 2 5

2

1/2

= (1 – 2 )

2 (1 – ) +

128 (1 – )81

p na n

p na

ÈÎÍ

˘˚

ÈÎÍ

˘˚˙

(15.29)

The analysis showed that a material containing short cracks, much smallerthan a characteristic length Lm, would propagate in an unstable manner atDTc, due to the released elastic energy being converted into kinetic energy,towards a final crack length Lf and cause a drastic reduction in strength. Forinitially longer cracks, i.e. L > Lm, or for short cracks that have reached Lf,

Table 15.1 Values of the thermal shock resistance parameters R, R¢, R¢¢¢¢ for a rangeof ceramic materials where HPSN is hot pressed silicon nitride and RBSN isreaction bonded silicon nitride (reprinted from Table 11.1 on p 213 of ‘Ceramics:Mechanical Properties, Failure Behaviour, Materials Selection’ by Munz and Fett,1999, published with permission from Springer-Verlag GmbH)

Al2O3 MgO ZrO2 SiC Bi3N4 BeO Al2TiO5

HPSN RBSN

R (K) 73 46 324 206 495 342 47 962R ¢(kW m–1) 2.19 3.9 2.7 66 75 20 36 9.6R≤¢¢(mm) 0.23 0.28 0.11 0.12 0.11 0.10 0.71


quenching at higher values of DT causes propagation in a stable, quasi-staticmanner and the material shows a gradual strength decrease above the DTc

associated with that particular crack size. In addition, it was shown thatthermal shock resistance increased with increasing initial microcrack density.

Evans (1975), Evans and Charles (1977), and Emery (1980) performedmore refined fracture mechanics studies regarding the onset and arrestconditions; Bahr et al. (1988) and Pompe (1993) extended this work andconsidered the propagation of multiple cracks; while Swain (1990) foundthat materials showing non-linear deformation and R-curve behaviour havea better resistance to thermal shock. More specifically, the behaviour of acrack in the thermal shock-induced stress field was deduced from thedependence of the crack length on the stress intensity factor. Unstablepropagation of a flaw in a brittle material under conditions of thermal shockwas assumed to occur when the following criteria were satisfied:

K KdKdL

dKdL

> , > ii (15.30)

where K is the thermal stress intensity factor, Ki is the material fracturetoughness (Ki) or the crack length-dependent critical stress intensity factor(KR) for materials that exhibit R-curve behaviour, and a is the crack length.In the case where

K KdKdL

dKdL

> , < ii

(15.31)

the flaw will propagate in a stable manner. Thus, by superimposing the Ki-curve of a material onto curves that describe the K-curve behaviour generatedfor a given thermal shock treatment, conditions of crack propagation andarrest were predicted. Such analyses verified Hasselman’s findings but wereable to define onset and arrest conditions with better accuracy.

The analyses were also found to correlate well with experimental findings.High-performance engineering ceramics usually have high stiffnesses combinedwith low values of toughness. Due to careful processing conditions the numberand size of flaws they contain are limited, ensuring high strengths. Theextent of the temperature differentials that they can sustain without crackingis then dictated mainly by the values of CTE and, to a lesser degree, by thevalues of thermal conductivity. Materials with lower CTE and high thermalconductivities can sustain higher values of DT. However, all such materialssuffer large and abrupt losses of strength at DTc as crack propagation occursin an unstable fashion. By contrast, refractory or porous ceramic materialsusually have low stiffnesses and contain a lot of large flaws or pores. Thesematerials do not show a definite DTc but exhibit a gradual reduction instrength starting at low values of DT. Examples of both types of materials aregiven in Fig. 15.4.


15.5 Thermal shock of particle- and

whisker-reinforced CMCs

The attraction of reinforcing ceramic matrices with particles or whiskers isthat, with appropriate microstructural design and property tailoring, materialswith property combinations not possible in monolithic ceramics can be obtained.In addition, the materials remain effectively isotropic and can be manufacturedby well-established techniques already in use for the manufacture of monolithicceramics (Hansson and Warren, 2000).

DTc DTc¢

0 200 400 600 800DT(∞C)

(a)

400

200

0

Str

eng

th (

MP

a)

Sintering temperature2100∞C2000∞C1900∞C

0 200 400 600 800DT(∞C)

(b)

50

40

30

20

10

0

Str

eng

th (

MP

a)

15.4 The thermal shock behaviour of (a) monolithic alumina, and (b)porous SiC at different sintering temperatures (reprinted from Figure11.6 on p 212 of ‘Ceramics: Mechanical Properties, Failure Behaviour,Materials Selection’ by Munz and Fett, 1999, published withpermission from Springer-Verlag GmbH.


It is the capability of property tailoring that gives particle- and whisker-reinforced CMCs the edge over monolithic ceramics under conditions ofthermal shock. By choosing carefully the properties of the reinforcement,reductions in Young’s modulus and CTE combined with increases in thermalconductivity compared with the unreinforced matrix material can be realised.Strict microstructural control during processing can result in fully dense,finely grained materials with good adhesion between reinforcement and matrixthat ensure high strengths. In this way, high critical temperature differentialsfor crack initiation can be achieved. In addition, the presence of thereinforcement results in the introduction of a number of energy-dissipatingmechanisms such as crack deflection, crack bridging, etc., which significantlyimprove toughness and damage tolerance. Thus, better resistance to crackpropagation compared with monolithic ceramics is also possible. The resultis a material that can sustain higher values of DT and, in addition, retain ahigher percentage of its initial strength at DT > DTc compared with its monolithicequivalent.

A number of experimental studies support the above analysis. Aghajanianet al. (1989) reported that the DTc of low-porosity alumina–matrix CMCsreinforced with aluminium particles increased compared with unreinforcedalumina, while porous CMCs with the same constituents behaved as refractoryceramics, i.e. displayed a low, but not definite, DTc and gradual reduction instrength with increasing DT. Similar refractory-type behaviour was observedby Aldridge and Yeomans (1999) in the case of a sintered alumina–matrixcomposite reinforced with 20 vol% iron particles that contained increasedlevels of porosity. However, a similar hot-pressed CMC with low porosityexhibited much higher DTc and higher strength retention at DTc comparedwith monolithic alumina (Fig. 15.5). Jin and Batra (1999) showed theoreticallythat crack bridging by metal particles resulted in a significant reduction ofthe thermal shock-induced stress intensity factor.

Bannister and Swain (1990) and Swain (1991) investigated the thermalshock behaviour of ZrO2-particle-reinforced Al2O3 and AlN- and BN-particle-reinforced TiB2 and reported higher thermal shock resistance compared withthe respective monoliths as well as no significant reduction in post-shockflexural strength. This was attributed to the materials exhibiting R-curvebehaviour due to the formation of microcracks around the reinforcing phases.Wang et al. (2001) found improved resistance to thermal shock (by ~70∞C)of a 6 vol% tungsten carbide particle reinforced alumina compared with theunreinforced material, which was consistent with higher toughness, reductionsin Young’s modulus and CTE, as well as the strong bonding of the reinforcementparticles to the matrix (so that they did not act as strength-reducing flaws).Similar reasons were put forward by Uribe and Baudin (2003) to explain theincreased thermal shock resistance of an alumina–matrix CMC reinforcedwith 10 vol% aluminium titanate particles, and also by Nieto et al. (2004),


who observed increased resistance to crack initiation and stable crackpropagation in an alumina containing 10 vol% sub-micron-sized AlN particles.In the case of alumina/silicon carbide-particle nanocomposites, Maensiriand Roberts (2002) observed superior resistance to thermal shock comparedwith the matrix material. However, since no changes in thermal or mechanicalproperties could be identified, the improvement was associated with a changein crack path (intergranular in alumina, transgranular in the nanocomposite).

Similar observations have been made regarding the thermal shock behaviourof whisker-reinforced CMCs. Tiegs and Becher (1987) reported no decreasein flexural strength following quenches into boiling water for a 20 vol% SiCwhisker-reinforced Al2O3. The authors noted that a small increase in thermalconductivity and a slight decrease in CTE of the composite, compared withthe matrix material, could not account for the extent of the improvement inthermal shock resistance and attributed it to the interaction of microcrackswith the reinforcement (i.e. crack arrest, deflection, etc.), which resulted inincreased toughness. Similar conclusions were drawn by Collin and Rowcliffe(2001) who tested the same material using the indentation-quench method.Pettersson and Johnson (2003) identified a clear correlation between increasedtoughness and pronounced R-curve behaviour with improved thermal shockresistance to explain the behaviour of alumina reinforced with Ti (C, N)

Hot-pressed Al2O3

Hot-pressed Al2O3 – FeSintered Al2O3 – Fe

0 100 200 300 400 500 600 700 800Temperature differential (∞C)

800

700

600

500

400

300

200

100

0

Ret

ain

ed s

tren

gth

(M

Pa)

15.5 The thermal shock behaviour of hot-pressed alumina, hot-pressed and sintered alumina reinforced with iron particles (reprintedfrom Journal of the European Ceramic Society, 19, Aldridge andYeomans, ‘The thermal shock behaviour of ductile particle toughenedalumina composites’, 1769–1775, copyright 1999, with permissionfrom Elsevier).


whiskers. Zhao et al. (2002) investigated the thermal shock resistance of b–sialon matrix composites reinforced with titanium carbonitride whiskers andnoticed that the addition of whiskers had no influence on the matrixmicrostructure, but their presence improved both the hardness and the fracturetoughness of the CMCs. No unstable crack extension occurred in the compositesfor DT = 90–700∞C, but above 700∞C performance deteriorated as a result ofsevere oxidation of the whiskers.

Studies have also shown that, since the amount of reinforcement addedaffects all mechanical and thermal properties, there is an optimum volumefraction of particle or whisker reinforcement that should be added to thematrix material to ensure superior resistance to thermal shock (Becher, 1981;Jia et al., 1996; Sbaizero and Pezzotti, 2003; Pettersson and Johnson, 2003).The shape of the reinforcement also plays an important role in determiningbehaviour under thermal shock. Sbaizero and Pezzotti (2003) showed thatthe use of coarse and elongated particles resulted in better CMC performancecompared with the use of fine-grained particles.

The importance of careful tailoring of the constituents to achieve improvedthermal shock resistance is highlighted by the study of Jia et al. (1996). Theincorporation of SiC whiskers in Si3N4 resulted in the CMC having a lowerDTc than monolithic Si3N4, as the CMC had a slightly lower thermalconductivity but a much larger CTE compared with the unreinforced matrix.However, stable crack growth occurred in the CMC in contrast to unstablecrack growth in the monolithic material, which was attributed to the presenceof the reinforcement. It was concluded that unreinforced Si3N4 is more suitablefor use under mild thermal shock conditions, where the objective is to avoidfracture, while the CMC should be used under severe thermal shock conditions,where initiation of cracking is unavoidable and resistance to crack propagationand post-shock strength retention become important.

The behaviour of particle- and whisker-reinforced CMCs under conditionsof thermal shock can be modelled successfully using the fracture mechanicsmethods outlined in the previous paragraph (e.g. Aldridge and Yeomans,2001) while the thermal shock parameters (figures-of-merit) can also beuseful for initial material comparison.

15.6 Thermal shock of fibre-reinforced CMCs

15.6.1 Introduction

Although particle- and whisker-reinforced CMCs can exhibit better thermalshock behaviour compared with monolithic ceramics, generally they stillshow a step decrease in their strength at DTc. A combination of the propertiesof high-performance engineering ceramics with high DTc and gradual strengthreduction above DTc (i.e. refractory-type behaviour) can only be realised


with the incorporation of continuous ceramic fibres into ceramic matrices.With optimum selection of fibres and matrices, favourable residual stress

conditions can be established in the matrix, which lead to increased DTc.Above DTc, matrix cracks appear but the presence of crack-deflecting fibre-matrix interfaces ensures minimal effect on mechanical properties as thefibres remain largely unaffected. As damage is also confined mostly to thesurface of the materials, changes in mechanical and thermal properties aremore readily identified by means other than mechanical testing.

In the following paragraphs an overview of damage due to thermal shockand its effect on the mechanical properties of CMCs with different fibrearchitectures is provided for a number of different reinforcement architectures.Subsequently, the effect of thermal shock on interfacial properties is discussed,followed by a description of attempts to analyse and model the thermalshock behaviour of these materials.

15.6.2 Thermal shock damage and its effect onmechanical and thermal properties

Unidirectional (UD) CMCs

Bhatt and Phillips (1990) reported that thermal shock reduced the flexuralmechanical properties of a UD composite comprising SiC fibres in a reaction-bonded Si3N4 matrix but that it did not affect its tensile properties (Young’smodulus, ultimate strength, matrix cracking stress). It was suggested that theloss in flexural strength was caused by the loss of inter-ply integrity of thecomposite after matrix fracture and the failure mode changing from a tensilefracture to delamination driven by shear stress.

Matrix cracking due to thermal shock and its effect on the flexural propertiesof UD Nicalon™ fibre-reinforced composites with borosilicate glass (Pyrex™)and lithium aluminosilicate (LAS) matrices was described by Kagawa et al.(1989, 1993). Damage was confined to the surface (two to three fibre diametersdeep) and was independent of DT. The Pyrex™–matrix system exhibitedmultiple matrix cracking perpendicular to the fibre axis at DTc = 600∞C,which coincided with a notable decrease in Young’s modulus E and flexurestrength. The decrease in E was attributed to matrix crack formation on thespecimen surface, but the reduction in flexure strength was explained as achange in failure mode to interlaminar shear failure, caused by a reduction ininterfacial shear strength due to thermal shock.

In the LAS–matrix system matrix cracks parallel to the fibre axis wereobserved at DTc = 800∞C, accompanied by a reduction in Young’s modulus,although flexure strength seemed to remain unaffected by thermal shocktreatment. This was attributed to the difference in the direction of matrixcrack propagation in the two composites due to the formation of a-spodumene–


silica solution in the LAS matrix during thermal shock, which could haveacted as a source of microcracking because of thermal expansion mismatch.

Multiple matrix cracking perpendicular to the fibre axis was also reportedby Blissett et al. (1997) for a UD Nicalon™/CAS (calcium aluminosilicate)(Fig. 15.6). The density of these cracks increased with increasing DT butshowed a reduction for DT > 800∞C, which seemed to be consistent with theformation of strong silica bridging between the matrix and the fibres.

In addition to matrix cracking perpendicular to the fibre axis, matrixcracks also occurred parallel to the mid-plane of the laminate. These crackswere first seen on the end faces of the composite at DTc = 400∞C (Fig. 15.7).The depth the cracks penetrated into the matrix increased and their pathgeometries changed with increasing DT. These effects were attributed to theinteraction of increasing applied thermal stresses with simultaneous reductionsin the interfacial shear strength due to oxidation of carbon. Similar damagemodes, termed ‘thermal debond cracks’, were observed by Graham et al.(2003) on the end face of a thermally shocked UD Nicalon™/LAS II composite.The authors highlighted the presence of high tensile radial stresses across thefibre–matrix interface, which favoured the appearance of such cracks, andnoted that they tended to run through fibre-rich regions where these stressesare highest. Reductions in thermal diffusivity due to thermal debond crackformation were measured. The appearance of such damage modes may also

20 mm

15.6 Multiple matrix cracking perpendicular to the fibre axis due tothermal shock in UD Nicalon™/CAS (reprinted from Journal ofMaterials Science 32(2) 1997, ‘Thermal shock behaviour ofunidirectional silicon carbide reinforced calcium aluminosilicate’Blissett, Smith and Yeomans, Figure 2, with kind permission ofSpringer Science and Business Media).


be responsible for the change in failure mode under flexure reported forother systems (Bhatt and Phillips, 1990; Kagawa et al., 1993).

Blissett et al. (1998) reported that thermal shock effects on the residualflexural properties of the Nicalon™/CAS were more evident at intermediatetemperature differentials, i.e. DT = 450–600∞C, and this was attributed to theobserved matrix cracking.

50 mm

(a)

(b)

15.7 (a) Photomicrograph, and (b) schematic of matrix cracking onend face of UD Nicalon™/CAS (reprinted from Journal of MaterialsScience 32(2) 1997, ‘Thermal shock behaviour of unidirectionalsilicon carbide reinforced calcium aluminosilicate’, Blissett, Smithand Yeoman, Figure 1a, with kind permission of Springer Scienceand Business Media).


Cross-ply CMCs

Blissett (1995) studied cross-ply Nicalon™/CAS laminates of two differentlay-ups and reported that thermal shock damage within individual plies wassimilar to that seen in UD specimens of the same material (Blissett et al.,1997).

Initial damage in a [0 /90 ]2 4 s∞ ∞ laminate was sustained at DTc = 400∞C in

the eight central 90∞ plies, and consisted of a single thermal debond cracksimilar to the ones observed on the end faces of UD Nicalon™/CAS (Fig.15.8). As this damage mode is not observed under monotonic tensile orfatigue loading applied along the axis of the longitudinal fibres in the 0∞plies (e.g. Pryce and Smith, 1992), its appearance is indicative of the biaxialnature of the thermal shock-induced stress field. Short cracks just crossingthe interface between plies as a result of thermal shock treatment were alsoreported. It was noted that both damage modes became more pronounced athigher values of DT.

The second laminate, [0∞/90∞]3s, exhibited only slightly different crackingfeatures, attributed to the difference in the stacking sequences of the laminates.A major thermal debond crack appeared at DTc = 350∞C and was confined tothe two central 90∞ plies. Similar cracks were observed in some of theadjacent 90∞ plies but were less pronounced. At DT = 400∞C matrix cracksperpendicular to the longitudinal fibres appeared in the 0∞ plies. For highervalues of DT, debond cracks were observed in most of the other 90∞ plieswhile the perpendicular matrix cracks were seen crossing to the adjacent 90∞plies before being arrested by the horizontally running debond cracks. However,the outer plies and the thinner 0∞ plies seemed to remain intact up to DT =

500 mm

15.8 Photomicrograph of thermal debond crack in the eight central90∞ plies of a [0 / 90 ]2 4 s

∞ ∞ , Nicalon™/CAS laminate (Blissett, 1995,reprinted courtesy of Dr M.J Blissett).


720∞C. Flexure strength and Young’s modulus were found to decrease withincreasing DT (Blissett et al., 1998).

2-D and 3-D woven CMCs

CMCs with 2-D woven fibre reinforcements have been found to possesshigher resistance to thermal shock than unidirectional or cross-ply CMCs ofthe same constituents (Nicalon™ fibres and SiC matrices) and prepared bythe same method (Chemical Vapour Infiltration-CVI) (Wang et al., 1997).

Only a slight drop in the flexural strength of a woven Nicalon™/Al2O3

composite was observed by Fareed et al. (1990) after quenching through DT= 1000∞C and 1200∞C. This was attributed to the effectively engineeredweak fibre/matrix interface.

Lamicq et al. (1986) reported that the bending strength of water-quenchedwoven SiC/SiC (CVI) specimens decreased slightly in the quench range DT= 300–750∞C, and then remained unchanged up to DT = 1200∞C. The compositealso seemed to exhibit a steep R-curve behaviour.

Wang et al. (1994, 1996) reported on the thermal shock behaviour of 2-Dwoven Nicalon™/SiC CMCs manufactured by CVI and polymer impregnationand pyrolysis (PIP), as well as that of a Nextel™–312/SiC (CVI) compositesystem. The Nextel™/SiC (CVI) system failed in post-quench flexure testsby fracture through the 2-D fibre planes and showed different criticaltemperature differentials for the onset of decrease in each of its macroscopicproperties. Reduction in ultimate strength, su, began at DTc(su) = 400∞C,matrix cracking stress, smc, started to decrease at DTc(smc) = 600∞C, whilethe work of fracture (WOF) decreased continuously as DT increased. Reductionsin thermal diffusivity with increasing values of DT were also reported for thissystem by Ellingson (1995).

The properties of the Nicalon‰/SiC (PIP) system followed a similar pattern(DTc(su) = 400∞C, DTc(smc) = 500∞C), though this system failed through aninterlaminar shear failure process (delamination) and the property reductionsaturated at DT = 600∞C. The Nicalon™/SiC (CVI) system failed by fracturethrough fibre planes but its properties (su, smc, WOF) had the same criticaltemperature difference, DTc = 700∞C. The pre- and post-quench stress–displacement curves for this material can be seen in Fig. 15.9. However,measurement of the Young’s modulus of this system before and after quenchingby means of a dynamic mechanical resonance technique showed the onset ofdecrease at DTc(E) = 400∞C, i.e. significantly lower than the DTc of the otherproperties.

An assessment of the thermal shock damage of woven Nicalon™/SiC(CVI) composite specimens was performed by Webb et al. (1996), the resultsbeing confirmed subsequently by Kagawa (1997). It was noted that the manypores and irregularities in the matrix inherent to this particular composite


geometry provide stress concentrators that amplify the thermal loading andcreate preferential sites for crack formation. For this reason, CVI-SiCcomposites exhibit lower DTc for the onset of cracking than monolithic SiC(Kagawa, 1997). Three types of thermal shock-induced damage on the materialsurface were reported:

∑ Matrix cracks that originated from the corners of uninfiltrated pores inregions outside fibre bundles. These cracks appeared at DT = 250∞C anddid not penetrate deeply into the fibre bundles, though the penetrationdepth increased with increasing DT.

∑ Matrix cracks between fibres within a fibre bundle. These occurred at DT= 1000∞C and were similar to thermal shock damage observed by Kagawaet al. (1993) in UD CMCs.

∑ Degradation of the fibre–matrix interface and removal of fibres. This typeof damage appeared at DT = 600∞C but was attributed to both thermalshock and/or oxidation effects.

In addition, matrix cracks that severed ligaments between cloths wereseen at DT ≥ 600∞C in the interior of thermally shocked specimens. Themechanism of formation of these cracks is not clear as thermal shock loadinginduces mainly high stresses at or near the surface. However, Kastritseas etal. (2004a) observed such cracks on polished parallel surfaces of similarSiC/SiC CMCs as well. Webb et al. (1996) reported that further increases inDT increased the severity of all types of thermal shock damage.

0 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 1.8 2Displacement (mm)

600

500

400

300

200

100

0

Str

ess

(MP

a)

15.9 Effect of increasing DT on stress–displacement curves ofNicalon™/SiC (CVI) – solid line corresponds to unshocked sample(reprinted from Wang et al. 1996, ‘Thermal shock behaviour oftwo-dimensional woven fiber-reinforced ceramic composites’,Journal of the American Ceramic Society, with kind permission ofBlackwell Publishing).


Correlation of these observations with property measurements by Wanget al. (1996) led to the postulation that surface matrix cracks that appear atlow DT (= 250∞C) are not strength-controlling but are responsible for thereduction in Young’s modulus observed at DTc(E) = 400∞C. On the otherhand, the interior cracks that severed links between fibre cloths at DT ≥600∞C seem to affect the strength of the composite, which decreases afterDTc = 700∞C. Such behaviour was summarised by Boccaccini (1998) in thegraph of Fig. 15.10. It has to be noted that the behaviour of E is also typicalof some thermal properties of CMCs (Ellingson, 1995; Graham et al., 2003).Note that there is no abrupt change in any property above DTc. However, ifthe fibre–matrix interface is strong, fibre-reinforced CMCs revert to behaviourtypical of monolithic ceramics (Twitty et al., 1995).

Damage modes resulting from thermal shock were identified by Kastritseaset al. (2004b) for a woven Nicalon™/CAS. The surface of the CMC wasdescribed as an assembly of alternating ‘plies’ containing either longitudinalor transverse fibres, with matrix-rich regions in between. Multiple matrixcracks confined to the surface appeared perpendicular to the fibre directionat DTc for the plies with longitudinal fibres and for the matrix-rich regions,while debond cracks running parallel to the longitudinal fibres could be seenin the plies with transverse fibres located towards the centre of the materialface. With increasing values of DT, debond cracks grew significantly inlength and depth (Fig. 15.11) but their number did not change significantly.By contrast, perpendicular matrix cracks did not change in morphology (i.e.they remained surface features of small depth) but increased moderately inlength and significantly in number. Crimp regions were not observed toaffect either crack initiation or development.

Pro

per

ty

DTDTc

s

E

Q–1

e

15.10 Schematic diagram showing the variation of fracture strength(s), Young’s modulus (E), internal friction (Q–1), and microcrackingdensity (e) with increasing shock severity (reprinted from ScriptaMaterialia, 38, Boccaccini, ‘Predicting the thermal shock resistance offibre reinforced brittle matrix composites’, 1211–1217, copyright1998, with permission from Elsevier).


The thermal shock behaviour of a 3-D carbon fibre-reinforced SiC-matrixCMC manufactured by CVI was assessed using the air-quench method byYin et al. (2002). Damage consisted of matrix cracks that induced a reductionin Young’s modulus, strength, and work of fracture for DT > 700∞C.

Studies of the interface

It appears that the strength of the fibre–matrix bond in thermally shockedCMCs remains unaffected unless high-temperature oxidation processes areinvolved. Boccaccini et al. (1999) did not identify any significant changes inthe properties of the fibre–matrix interface of SiC/borosilicate glass compositesas a result of thermal shock, while Chawla et al. (2001) observed only aslight decrease in the interfacial shear stress of a thermally shocked Nicalon™-fibre SiC-whisker BMAS (barium magnesium aluminosilicate)–matrix hybridcomposite. If heating and soaking at temperatures harmful to the integrity ofthe interface are involved prior to quenching, degradation due to oxidationprocesses occurs (Blissett et al., 1997, 1998). In this case, changes in propertiesare explained as a combination of both oxidation and thermal shock (Grahamet al., 2003). The oxidation of the carbon interface in Nicalon™-reinforcedglass ceramic–matrix CMCs leads to cracks in the matrix, causing fibre

15.11 Matrix cracking due to thermal shock in a 90∞ ply of a wovenNicalon™/CAS at (a) DT = 700∞C, and (b) DT = 800∞C (after Kastritseaset al, 2004b).

(a)

(b)


failure due to the resulting strong interfacial bond (Blissett et al., 1997;Kastritseas et al., 2004b).

15.6.3 Theoretical considerations

Only a few studies have appeared in the literature regarding the analysis andmodelling of the thermal shock behaviour of fibre-reinforced CMCs.

Wang and Chou (1991) studied numerically the 3-D transient thermalstress in angle-ply laminated composites caused by sudden changes in thethermal boundary conditions. The study showed that DTc would be reducedif the fibre volume fraction, Vf, the CTE or the Young’s modulus of thecomposite increased, while it would increase with increasing thermalconductivity. By contrast, Boccaccini (1998) showed that increasing Vf

increased the DT for the onset of matrix cracking in glass and glass–ceramicmatrix composites. Wang and Chou (1991) also demonstrated that the changein CTE had the biggest effect on DTc while the change in thermal conductivityhad the least influence. In addition, as the fibre orientation angle deviatesfrom 45∞ towards 90∞ or 0∞ the interlaminar normal stress decreased whilethe in-plane thermal stress transverse to the fibre direction increased. Thisresulted in the initial failure mechanism changing from delamination to matrixmicro-cracking.

Wang et al. (1996) performed a 1-D qualitative analysis using the stressesgenerated due to thermal shock and the residual stresses associated with thethermal expansion mismatch between the fibres and the matrix. The analysisshowed that if (CTE)f > (CTE)m then the matrix is under tension only in theradial direction and possible matrix cracking will be circumferential, whilethe fibre is under tension in all directions (longitudinal, radial, andcircumferential), which may promote fibre damage. If (CTE)f < (CTE)m,then the matrix is under tension in both longitudinal and circumferentialdirections; hence, radial and normal-to-fibre matrix cracking will be possible.Moreover, the fibre is under compression in all three directions, so fibredamage will be limited. Debonding would also be possible in both cases. Asa confirmation, they applied their analysis to the results of Kagawa et al.(1993). In the Nicalon™/Pyrex™ composite, where (CTE)f < (CTE)m, thematrix is under a tensile stress in the longitudinal direction, which dictatesthat cracks will be perpendicular to the fibre axis, as observed in the experiment.Conversely, in the LAS–matrix composite (CTE)f > (CTE)m, i.e. the matrixis under tension in the radial direction, which results in cracks parallel to thefibre.

Particular interest has been paid to the analytical prediction of the DTc forthe onset of matrix cracking. Blissett et al. (1997) and Boccaccini (1998)considered the residual stresses present in the composite due to thermalexpansion mismatch between fibre and matrix which, when superimposed to


the applied thermal stresses, could lead to matrix cracking. Their approachwas based on the assumption that the stress that produces matrix crackingwould be the same whether applied mechanically or thermally. Hence, thematrix cracking stress (smu) was equated with the critical thermal shock-induced stress (sTS), which is the thermal stress required to produce matrixcracking, taking also into account the effect of residual stress (sr), i.e.

s s smu cTS

r = + (15.32)

Following equation (15.18), the critical thermal shock-induced stress is givenas:

s anc

TS c = 1 –

AE TD(15.33)

For Blissett et al. (1997), E, a, and v are matrix properties whereas Boccaccini(1998) defines ‘effective’ values calculated using the rule of mixtures.

Blissett et al. (1997) used the concentric cylinder model of Powell et al.(1993) to obtain residual stresses, whereas Boccaccini (1998) utilised theresults of a simple force balance in 1-D performed by Wang et al. (1996),which gives the residual thermal stresses in the matrix along the axial directionas:

s armatrix m

m f

f f

= 1 +

(1 – )E TE V

E V

D DÏÌÓ

¸˛

(15.34)

Different models were also used to obtain the matrix cracking stress (smu),with Blissett et al. (1997) using the classic Aveston et al. (1971) (ACK)analysis and Boccaccini (1998) using the model of Pagano and Kim (1994),which gives smu as:

s muIC =

2 +

K

r sp

(15.35)

where KIC is the fracture toughness of the matrix, r is the fibre radius, and sis the fibre spacing. The model assumes that there is no interaction betweencracks, which Boccaccini (1998) explains as a plausible assumption in theearly stages of thermal shock damage.

By solving the resulting expressions for DTc, the values of the criticaltemperature differentials are obtained. These are given as:

DG

T vAE

E E VE rVc

m m

m f m f2

1 m

13

r = 1 – 6 – a

ts

ÊËÁ

ˆ¯

È

Î

ÍÍÍ

˘

˚

˙˙˙

(Blissett et al., 1997)

(15.36)


D D DT

vAE

K

r s

E TE V

E V

e

e ec

IC,m m

m f

f f

= 1 –

2 + –

1 + (1 – )a

p

aÈ

Î

ÍÍÍ

˘

˚

˙˙˙ (Boccaccini, 1998)

(15.37)

In (15.36) Gm is the matrix fracture energy, t is the interfacial shear strength,and E1 is the axial modulus of the composite. In (15.37) e refers to theeffective properties of the composite, which, for unidirectional fiberreinforcement, can be calculated with good approximation by the rule ofmixtures.

Although the two approaches are very similar, the value of DTc inBoccaccini’s model does not depend on the interfacial shear strength t, as aresult of the model chosen for the value of matrix cracking stress. Blissett etal. (1997) suggested that their method was valid for the UD material providingthat some key parameters (interfacial shear stress, matrix fracture energy)were determined independently.

A recent analysis by Kastritseas et al. (2004c) suggested that in both casesthe magnitude of the thermal shock-induced stresses was overestimated asthe anisotropic character of the materials was not taken into account. Ifmaterial anisotropy is accounted for, then both (15.36) and (15.37) cannotpredict DTc accurately even for the largest possible value of the thermalshock-induced stresses (corresponding to a maximum value of the stressreduction factor, A = 0.66). To explain the discrepancy, it was proposed thatthe interfacial properties may be affected by the shock due to the biaxialnature of the induced stress field, which dictates that a tensile thermal stresscomponent that acts perpendicular to the fibre–matrix interface is present forthe duration of the shock.


This chapter has reviewed the performance of CMCs under conditions ofthermal shock. It has been shown that CMCs exhibit superior resistance tothermal shock, compared with their monolithic counterparts, as catastrophicfailure can always be avoided. Resistance to higher temperature differentialsand property retention after the onset of thermal shock cracking (especiallyin fibre-reinforced CMCs) can be realised, provided that the mechanical andthermal properties of CMCs are optimised by careful choice of theirconstituents.

The behaviour of particle- and whisker-reinforced CMCs can be adequatelydescribed by using and adapting the models and methodology developed formonolithic ceramics. By contrast, analysis and modelling of the performance


of fibre-reinforced CMCs is a subject still in its infancy that requires furtherattention. The situation is very complex due to the variety of damagemechanisms developed in these materials (especially 2-D CMCs) and isfurther complicated due to their anisotropic character, the scarcity ofexperimental results, and the variety of manufacturing methods that result inmaterials with different design philosophies.

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Wang, L., Shi, J.-L., Gao, J.-H., Yan, D.-S. (2001), ‘Influence of tungsten carbide particleson resistance of alumina matrix ceramics to thermal shock’, J. Eur. Ceram. Soc., 21,1213–1217.

Wang, Y.R., Chou, T.-W. (1991), ‘Thermal shock resistance of laminated ceramic matrixcomposites’, J. Mater. Sci., 26, 2961–2966.

Webb, J.E., Singh, R.N., Lowden, R.A. (1996), ‘Thermal shock damage in a two-dimensionalwoven-fiber-reinforced-CVI SiC-matrix composite’, J. Am. Ceram. Soc., 79(11), 2857–2864.

Wereszczak, A.A., Scheidt, R.A., Ferber, M.K., Breder, K. (1999), ‘Probabilistic thermalshock testing using infrared imaging’, J. Am. Ceram. Soc., 82(12), 3605–3608.

Yin, X., Cheng, L., Zhang, L., Xu, Y. (2002), ‘Thermal shock behavior of 3-dimensionalC/SiC composite’, Carbon, 40, 905–910.

Zhao, Z., Johnson, M., Shen, Z. (2002), ‘Microstructure and mechanical properties oftitanium carbonitride whisker reinforced b–sialon composites’, Materials ResearchBulletin, 37, 1175–1187.

434

16.1 Introduction

Superplasticity is macroscopically defined as the ability of a polycrystallinematerial to exhibit large elongations at elevated temperatures and relativelylow stresses. It is commonly found in a wide range of materials from metalsto ceramics (bioceramics or high-temperature superconductors, among others)when the grain size is small enough: a few micrometres for metals and lessthan a micron in ceramics.

From a microscopic point of view, a superplastically deformed polycrystalis characterized by a microstructure, i.e., grain size and form factor, almostunchanged after deformation, and it is generally accepted that these featuresare mainly achieved through grain boundary sliding (GBS). The mechanismsaccounting for the relaxation of the stresses created during GBS (what iscommonly known as the accommodation process), especially in ceramics,are not yet clearly elucidated and remain controversial. However, it is acceptedthat atomic diffusion, either between the grain boundaries and the bulk, oralong grain boundaries or during the non-conservative movement ofdislocations, is among the accommodation processes controlling superplasticity.

Superplasticity is a very promising property, not only because, like inmetals, the superplastic formation opens a way for the manufacturing ofcomplex ceramic pieces for industrial applications, but also because thecombination of GBS and diffusional processes makes superplasticity aninteresting tool for joining ceramic pieces in shorter times and lowertemperatures than the diffusional joining technique.

Several parameters can influence strongly the superplastic behaviour ofceramics, i.e. the strain rate at which the material can be superplasticallydeformed. Between them can be mentioned the grain size, second phases andsegregation of impurities at the grain boundaries, etc.

This chapter discusses the following leading topics:

∑ Analysis of both the macroscopic and microscopic features of superplasticity

16Superplastic ceramic composites

A D O M Í N G U E Z - R O D R Í G U E ZD G Ó M E Z - G A R C Í A, Universidad de Sevilla, Spain

and F W A K A I, Tokyo Institute of Technology, Japan

Superplastic ceramic composites 435

∑ The different accommodation processes controlling superplastic output∑ A critical analysis of the parameters improving superplasticity∑ Applications: forming and joining∑ Future tendencies in the field, with a special emphasis on up-to-date

information about the most outstanding and promising superplasticity inrelated ceramic materials.

16.2 Macro- and microscopic superplastic

characteristics

Although superplasticity in metals has been extensively studied since the1960s, it was only in the 1980s when superplasticity in ceramics (monolithicand composites) started to become a very active research field, and it hasexpanded rapidly since then. The reason for this development has been theneed to improve the structural properties of these materials due to the demandsof industry, asking for materials to be used in more severe working conditions.Ceramics can be a good candidate if their mechanical behaviour, especiallyat low temperatures, can be improved. Several techniques have been developed;probably the most widely used is based on the transformation toughness inzirconia, a keystone in ceramics investigation since Garvie et al.1 publishedtheir paper ‘Ceramic steel?’. Recently, new techniques have emerged throughthe processing and sintering of ceramic powders, leading to the attainment offully dense ceramics with equiaxed grain sizes below 1 mm, primordial forsuperplastic behaviour.

The first observation of superplasticity in a 3 mol% yttria-stabilizedtetragonal zirconia polycrystal ceramic (YTZP) with a grain size of 0.4 mmwas reported by Wakai et al.2 in 1986 (Fig. 16.1). Since then, a large number

16.1 First demonstration of superplastic deformation of 3Y-TZP:specimens before and after deformation at 1450∞C and 3 ¥ 10–4 s–1.


of fine-grained polycrystalline ceramics and ceramic composites have beenshown to have superplastic behaviour.3–8

Today, from an engineering point of view, the name superplasticity isascribed to a polycrystalline material pulled out to very high tensile elongationsprior to failure with necking-free strain. This phenomenon is usually foundin many metals, alloys, intermetallics, composites and ceramics (recently inhigh-temperature superconductor ceramics) when the grain size is smallenough, less than 10 mm for metals and less than 1 mm for ceramics.

From a microscopic point of view, when a polycrystalline material isdeformed at high temperatures, grain boundary sliding (GBS) takes place intwo different ways:

∑ The deformation is due to the flow of point defects, then GBS occurs tomaintain grain coherency. This is called diffusional creep: Nabarro–Herringif the diffusion takes place along the bulk, or Coble if it takes place alonggrain boundaries. In these cases, each individual grain suffers almost thesame deformation as that imposed on the specimen and the grains whichare nearest neighbours remain nearest neighbours. This is termed ‘Lifshitzgrain boundary sliding’ (Fig. 16.2(a)).

∑ A different situation happens when GBS is responsible for the deformation.In order to release stresses created during GBS, deformation may beaccompanied by intergranular slip throughout adjacent grains, by localizedslip adjacent to the boundaries or by diffusional process of point defects.Sometimes, formation of triple-point folds or the opening up of cracks atthe triple points can also accommodate GBS; however, as soon as there iscoalescence of voids or cracks, the material fails, and that happens at lowstrains. This type of GBS, accommodated by the process described above,is termed ‘Rachinger grain boundary sliding’ (Fig. 16.2(b)).

When GBS is accommodated by some of the mechanisms involvingdislocation movement or diffusion of point defects, the grains retain almostthe original size and shape even after large deformations. This GBS, as theprimary mechanism for deformation, is the basis for the high ductility exhibited

s

s

s

s(a) (b)

16.2 (a) Schema of the Lifshitz GBS. The thick line shows the GBSafter deformation. (b) Schema of the Rachinger GBS, showing thatthe grain shape remains constant after deformation.


by some materials at high temperatures and therefore for their structuralsuperplastic behaviour.

Although superplasticity is defined as the ability of a polycrystalline materialto exhibit large elongations, in many ceramics-related materials and ceramiccomposites superplasticity is also said to occur even though the polycrystalis deformed in compression, or in three- or four-point bending conditions, aslong as GBS is the primary deformation process.4–7

The commonly used equation for the steady-state strain rate e characterizingsuperplastic behaviour is written as:

e s = A

GbkT

bd G

Dp n

ÊË

ˆ¯

ÊË

ˆ¯ (16.1)

where A is a dimensionless constant, G is the shear modulus, b is the magnitudeof a Burgers vector, k is the Boltzmann constant, T is the absolute temperature,d is the grain size, s is the stress, D = D0 exp(–Q/RT) is the appropriatediffusion coefficient involved in the accommodation process (D0 is the frequencyfactor, Q the activation energy and R the gas constant), and p and n are thegrain size and stress exponents respectively.

The values of the creep parameters (p, n and Q) identifying the superplasticbehaviour of ceramic-related materials are not unique to such materials, norto the same type of materials. As shown in the review papers, these parametersare very similar in tension as in compression in zirconia-based materials(probably the most widely studied ceramics in the widest experimentalconditions), although that depends strongly on the purity of the ceramics;5,7

however, their behaviour seems to be very different in compression than intension when an aid-sintering phase is necessary during the processing, as insilicon carbide and silicon nitride ceramics.8

For the sake of clarity, we will mention the reasons for the discrepanciesin the creep parameters in materials for which most studies have beenperformed. In the case of YTZP, values of p between 1 and 3, of n between2 and higher than 5, and of Q between 450 and 700 kJ/mol, have beenreported during creep.7

Three different explanations have been developed to account for thediscrepancies in the experimental creep data:

∑ These differences have been interpreted on the basis of two sequentialmechanisms: at high stresses, deformation occurs by GBS which is theslower process and therefore controls plasticity, whereas at low stressesthe deformation is controlled by an interface-reaction process. In bothcases the activation energy is the same.9,10 However, this explanation failsfor several reasons: the experimental values found for the activation energyare not constant (it is high at low stresses and becomes equal to 450 kJ/mol at high stresses) and the n values decrease gradually to 2 when thestress increases.2,9


∑ Another explanation was proposed by Berbon and Langdon11 using amodified Coble mechanism developed by Arzt et al.12 In this model, theyassume the existence of perfect grain-boundary dislocations evenly spacedin the boundary planes so that they can all climb at the same speed. Thismodel predicts that n values cannot be higher than 3 and the activationenergy is constant and equal to the energy for grain boundary diffusion ofZr. However, both predictions are contrary to the experimental observations.Finally, during Coble creep, grains of polycrystals change shape and sizeto reflect the overall strain within the sample, again contrary to themicrostructural observations of these deformed materials, in which grainsremain almost unchanged.

∑ Probably the most plausible explanation for the scatter of the creepparameters is based on a single mechanism involving GBS with a thresholdstress (s0).

7,10,13 When a threshold stress is introduced into the creep equation(16.1), all the creep parameters in YTZP become n = 2, p = 2 and Q = 460kJ/mol whatever the stress or temperature of the test. The value of this s0

was found experimentally:13

s 0–4 = 5 10

exp120kJ/mol

¥

ÊË

ˆ¯RT

d(16.2)

with d in mm.

Recently, a model has been developed on the basis of the segregation ofthe yttrium atoms at the grain boundaries to account for s0. These areresponsible for an electric field to appear, which influences the graindisplacement of each other (i.e. GBS). The model is able to explainquantitatively the dependence of s0 on both temperature and grain size.14

Another explanation for the threshold stress has been pointed out, whichtakes into consideration the stress required for intragranular dislocations tobe activated, estimated to be 3.6–10.7 MPa;15 however, the role theseintragranular dislocations play is not yet clear, because the flow stress requiredto activate dislocations in yttria-tetragonal zirconia single crystals attemperatures as high as 1400ºC is over 400 MPa, much higher than the stressused in superplasticity at these conditions.16

For second-phase sintered ceramics, these phases control the plasticityand they are responsible for the asymmetric behaviour when deformed intension or compression, because there is a crucial difference in themicrostructure evolution associated with tension and compression creep.There are few explanations for this asymmetry.

It is a well-established fact that under tension, the formation and nucleationof cavities contributes much more significantly to the macroscopic strainthan under compression.17 Experimental results have shown that the difference


between tensile and compressive creep can be approximately explained byconsideration of the volume fraction of cavities.18

Another explanation of the creep asymmetry is based on the rate of approach(during compression) or separation (during tension) of adjacent grain facetscontrolled by the viscous secondary phase and the fact that the number ofgrain facets supporting compressive stresses is less than those supportingtensile stresses.19

When the material is composed of a soft phase with rigid inclusions, thestrain rate is written in the form:

e s = (1 – )1 fA V q n (16.3)

where A1 is a stress-independent constant and Vf the volume fraction ofinclusions. The exponent q depends exclusively on the stress exponent, theaspect ratio of the rigid inclusions and the orientation with respect to thestress axis. This term has been proposed as the origin of the asymmetricbehaviour between tension and compression.20

None of these accommodation processes is diffusion-controlled, since adiffusional process cannot account for the asymmetric behaviour betweentension and compression, due to its inherent symmetry in the sign of thestress.

16.3 Accommodation processes controlling

superplasticity

As mentioned above, several mechanisms can be responsible for the grainboundary sliding accommodation; however, so far there is no consensus ona general single mechanism to accommodate GBS, nor one concerning aparticular ceramic. In this section the different mechanisms for accommodationwill be analysed. For the sake of clarity, the accommodation process will bedescribed for each type of ceramic, whether monolithic, with secondaryglassy phases or composite.

For monolithic ceramics and ceramic composites, during superplastic flow,the relative motion of two adjoining grains has components parallel andperpendicular to their common grain boundary. GBS is the component parallelto the grain boundary and is responsible for 70–80% of the deformation insuperplasticity of fine-grained polycrystals,21–23 as has been shown bymeasurements of the grain aspect ratio, by both SEM and atomic forcemicroscopy as displayed recently by Duclos24 in YTZP. During GBS, rigidgrains inevitably generate cavities and cracks, and in order to avoid fractureand allow the material to deform superplastically, accommodation processesinvolving non-conservative motion of the grains (movement perpendicularto the grain boundary) are necessary. This non-conservative motion involveseither diffusion of point defects, or dislocation glide and climb and grainboundary migration when grain growth occurs.


16.3.1 GBS accommodated by diffusional flow

Ashby and Verrall (thereafter A-V) developed a model for GBS accommodatedby a diffusional process which integrates quantitatively the essential topologicalfeatures found in superplasticity.25 The principle of this model for grainrearrangement is shown in Fig. 16.2(b). This type of grain rearrangementretains the equiaxed grains after large deformation as shown in the typicalmicrostructural features of the polycrystals superplastically deformed.7,24 Amodified A-V model accounting for a more realistic symmetrical diffusionpath was developed by Spingarn and Nix when diffusional flow occurs onlyalong grain boundaries.26 Several modifications of the original A-V modelhave been made by various authors.27 However, the main features of theoriginal A-V model remain unchanged.

The constitutive equation of the A-V model, when lattice and grain boundarydiffusion are taken into account, is written:

e s g d =

100 –

0.721 +

3.32 L

gb

L

WkTd d

DD

dDÊË

ˆ¯

ÊËÁ

ˆ¯

(16.4)

where W is the atomic volume of the diffusion controlling species, 72g /d isa threshold stress for the grain-switching event, g is the grain boundary freeenergy, d is the thickness of the boundary, and DL and Dgb are the lattice andgrain boundary diffusion coefficients, respectively.

As can be observed in Eq. (16.4), the strain rate is a linear function of thestress (n = 1) in disagreement with experimental observations for manymonolithic ceramics5,7,28 for which n is 2. For YTZP, the grain size dependencefits to 2 and the activation energy corresponds to that of cationic latticediffusion.

In the case of YTZP, on which a large number of studies have beenperformed, the data could be fitted to a constitutive equation, which is identicalto that found in metals when lattice diffusion is the rate-controllingmechanism:29

e s s = 3 10

( – ) exp –

460kJ mol10 02

2

–1

¥ ÊËÁ

ˆ¯Td RT

= 2 10 – 7 0

2 2

LZr¥ Ê

Ëˆ¯

ÊË

ˆ¯

GbkT G

bd

Ds s

(16.5)

16.3.2 GBS accommodated by dislocation movement

When GBS is accommodated by the movement of dislocations, the strainrate can be deduced in a similar way as for recovery creep: the dislocationsgenerated at the grain boundaries move either along the grain boundaries or


across the grain until they are piled up at an obstacle, which can be overcomewhen the stress at the head of the pile-up is high enough to promote theclimb of the head dislocation. From this picture the strain-rate-controllingGBS can be written:30

e s

2

3 3

2 3

2ª DL b

h GkTd(16.6)

with D an appropriate diffusion coefficient, L the length of the pile-up and hthe climb distance.

Depending on the grain size, two limit cases have been analysed.30 At lowgrain size, it happens that L ~ d, h ~ 0.3d and Eq. (16.6) becomes:

e s

2 gb2 2

ª ÊË

ˆ¯

ÊË

ˆ¯

A D GbkT

bd G

(16.7)

where A2 is a dimensionless constant with a value of 10.Conversely, when the grain size is large, L ª 20Gb/s, h ª bG/16s (for

more details see reference 30), and by substituting these parameters, Eq.(16.6) becomes:

e s 2 L3

ª ÊË

ˆ¯ÊË

ˆ¯

A D GbkT

bd G

(16.8)

where A2 is a dimensionless constant having a value of the order of 103.Recently, the activity of dislocations has been observed during the

superplasticity of YTZP.15 The authors suggest that the rate of deformationis controlled by the recovery of the intragranular dislocations in the high-stress region where n is 2.7 and p is between 2 and 3.

As can be inferred from the equations outlined above, none of the differentmodels can adjust the creep parameters for all the different ceramics, especiallyin the case of YTZP,7 explaining why there is still controversy over theaccommodation process controlling superplasticity. The same conclusionscan be outlined for ceramic composites, although more experimental workshould be done.20,31

For ceramics with secondary glassy phases, the accommodation processesare governed by these phases. Although diffusion may occur, the glassyphase viscosity controls accommodation mainly in different ways:

∑ These glassy phases may act as a lubricant for grain boundary sliding. Inthis case, the accommodation mechanism is the viscous motion of thesesecondary phases.

∑ The secondary glassy phases can improve the diffusivity pathwaysthroughout the grain boundaries. Accommodation is controlled by diffusionalong them.


∑ Finally, the secondary phase can provide a preferential location for thenucleation and growth of cavities during deformation, producing fractureof the material and reducing its superplasticity ability.

At high temperatures the glassy phase may become less viscous and evenliquid and as a consequence may account for the plastic deformation. However,viscous flow creep is not regarded as a viable creep mechanism forsuperplasticity due to its limited deformation, which corresponds to theredistribution of the glassy phase and therefore to the squeeze of these secondaryphases from grain boundaries subjected to compression.8

In the next section, we will analyse the mechanisms that are considered tocontrol superplasticity in ceramics with secondary phases.

16.3.3 Solution–precipitation creep

In this case, the secondary phases melt at temperatures lower than the matrixand, provided that the crystals are at least partially soluble in the glassyphase, creep may take place by:

∑ solution of the crystal in the liquid phase at grain boundaries undercompression,

∑ diffusion along the liquid phase, and∑ precipitation of the crystalline material at grain boundaries under traction.

The solution–precipitation creep model was first proposed by Raj andChyung;32 they assumed two cases:

∑ The strain rate is controlled by the diffusion along the glassy phase. Thus:

e h s = 1 13d

(16.9)

with h the viscosity of the secondary phase.∑ The strain rate is controlled by the reactions of solution and precipitation

at the interfaces. If this is the case:

e s = 1d

kd (16.10)

kd being a reaction constant.

This model postulates that the glassy phase in compression is able tosupport normal stresses because of the existence of islands, thus avoiding thecomplete squeeze of this intergranular liquid (Fig. 16.3). However, it hasbeen shown that these islands are not necessary for the grain boundaries tosupport normal stresses. Several modifications and revisions of this firstmodel have been made.8


An important modification of this model was performed by Wakai.33 Themain assumptions are that the solution and precipitation reactions take placeat line defects as ‘kinks’ in steps formed at the grain boundaries (Fig. 16.4),and the spacing between kinks is small enough for the step to be consideredas an ideal source or sink of solute particles. Thus, the solution and precipitationof crystalline materials at these steps produces their movement, andconsequently strain and strain rate will have an expression analogous toOrowan’s equation for dislocation movement:

e r = S Sav

d(16.11)

with rS the density of surface steps per unit length, a the height of the stepsand vS the velocity of the steps. This velocity depends on the process ofintegration into the crystal at a kink, the diffusion in the absorption layer and

h

Glassyphase

Island

Upper grain

Lower grain

Glassy phase

16.3 Schema of the solution–precipitation model with the islandstructure supporting normal stresses.

Grain

Grain

‘Kink’

Glassy phase

(i) Solute transport alongthe glassy phase

(iv) Integrationat a kink

(ii) Deposition at theadsorption layer

(iii) Diffusionalong this layer

16.4 Schema of the solution–precipitation in Wakai’s step model.


the diffusion in the liquid film. The rate-controlling process will be the moreresistant to the step movement.

Three different situations were analysed by Wakai for the density of steps:

∑ The constant density of surface steps is independent of the applied stress.∑ If the initial density of steps is very low, two-dimensional nucleation of

the surface steps occurs.∑ If the continuous source of steps is a screw dislocation, then a spiral step

is generated.

The combinations between the rate-controlling process and the density ofsteps give the different equations of Wakai’s model, which are summarizedin Table 5 of reference 8. The case of constant density of steps modelled byWakai is equivalent to the diffusion-controlled creep modelled by Raj andChyung32 and is also consistent in terms of the stress, temperature and grainsize dependence of the strain rate for interface-reaction-controlled creeppredicted by Raj and Chyung.32 However, in the two cases of bidimensionalnucleation of step and spiral step, the creep parameters differ from thosepredicted by wakai.33 In particular, for two-dimensional nucleation, there isa divergence of the creep parameters which has been recently solved34 byconsidering in detail the precipitation or solution of the crystalline materialat the step, which changes significantly the free enthalpy involved in theprocess. The essential key of this new modification is the free energy changeper unit volume for the precipitation mechanism, making a strong correctionto the apparent parameters n and Q measured in mechanical tests.

16.3.4 Shear thickening creep

This phenomenon has been postulated as an explanation for the compressivesuperplastic deformation of a SiAlON which undergoes a transition from n= 1 (Newtonian behaviour) to n = 0.5 for a characteristic critical stress(sc) independent of both temperature and composition of the secondaryphase.35

The model is based on the idea that the glassy phase is composed of twolayers – a normal glassy phase layer behaving in a Newtonian way, embeddedinto an over-condensed layer with non-Newtonian behaviour. Thus, for stresseslower than the critical stress, the creep is controlled by the normal glassyphase (n = 1), and when the stress exceeds a critical value, the squeeze of thisphase makes the two over-condensed layers come into contact, thus thematerial creeps in a non-Newtonian way (n = 0.5). The creep rate is written:

e sh = (1 – )f

2.5V ¢ (16.12)

with h¢ an apparent viscosity depending on temperature, grain size, phase


composition and liquid content, and Vf the volumetric fraction of the over-condensed rigid phase, which for s s 2

3 c≥ yields a value:

Vfc = 1 – 1

4 +

2ss (16.13)

explaining the transition from n = 1 to n = 0.5. Again a review of the modelcan be found in the literature.8

16.4 Parameters improving superplasticity

The strategy to enhance superplasticity is twofold: refinement of themicrostructure, i.e. decrease of the grain size, or improvement of theaccommodation process needed to relax the stresses created during GBSthroughout the appropriate diffusion coefficient. Although both means areindependent of each other, in a few cases the reduction of the grain size mayinduce a decrease in the diffusion coefficient involved in the accommodationprocess, giving rise to a compensating effect. These different strategies willbe analysed in the following sections.

16.4.1 Refinement of the microstructure

As observed in the different equations accounting for superplasticity in ceramic-based materials, the strain rate is an inverse function of the grain size; inconsequence, the grain size should be stable and as small as possible toattain high-strain-rate superplasticity. If grain growth occurs duringdeformation, the level of stresses for successive deformation will increase,inducing the formation of intergranular cavities leading to failure.

Several techniques have been developed to achieve ceramics and compositeswith fine microstructure:

∑ One of the techniques to suppress grain growth is based upon the inclusionof dispersed phases into the ceramics. With this regard, a multi-phaseceramic composite containing 40 vol% ZrO2, 30 vol% spinel and 30 vol%Al2O3 has been sintered,36 which was superplastically deformed to 1050%at 1650ºC at a strain rate of 0.4 s–1. Other zirconia-based composites havealso been fabricated with inhibition of grain growth.31,37 With this techniqueit is possible to deform the ceramics, with their microstructure unchanged,at a temperature at which grain growth would be important in ceramicswithout dispersed phases.

∑ Another technique makes use of ultra-fine powders sintered under stress-assisted conditions so the sintering temperature is reduced,38 for instancehot-isostatic pressing, hot pressing, sinter forging, or techniques with fast


heating and cooling ramp, such as spark plasma sintering (SPS) ormicrowave sintering, which avoid grain growth by reducing the time ortemperature of sintering. Very fast densification has been reported inoxides and Si3N4-based ceramics by SPS.39 Microwaves have also beenused to sinter PSZ40 and alumina.41

∑ The sintering of Y2O3 nanocrystalline ceramics (d = 60 mm) has beenachieved through a two-step sintering method. The first step is pre-sinteringat high temperatures to obtain ceramics with intermediate density valuesbetween 70 and 80%. Secondly, suppression of grain growth is achievedby sintering at lower temperatures than those used during the first step,exploiting the difference in kinetics between grain-boundary diffusion,which controls sintering, and grain-boundary migration, which controlsgrain growth (second step).42

16.4.2 Improvement of the processes controllingdiffusion in superplasticity

The diffusion coefficients of the process controlling superplasticity may beenhanced or retarded by the addition of impurities or solute atoms or by theaddition of secondary phases, normally used as sintering aids, which distributealong the grain boundaries and triple-point junctions of the grains.

There are a great number of papers dealing with the influence of the grainboundary segregation on superplasticity in YTZP. It has been shown that thesuperplastic flow stress at 1400ºC of a 3YTZP doped with different cationsis correlated with the ionic radius of the dopant.43,44 Cations with smallerionic sizes decrease the flow stress, whereas those with large ionic sizesincrease the flow stress. The authors suggest that the flow stress is determinedby the grain boundary diffusivity, which is affected by the segregation of thedopant. The same improvement of superplasticity has been respectivelyfound45,46 in a 0.3 mol% SiO2 doped 3YTZP with d = 0.35 mm and a 0.18mol% Al2O3 doped 3YTZP with d = 0.4 mm. This behaviour also seems toaccount for fine-grained Al2O3 with ZrO2 as a dopant.47,48

However, this explanation conflicts with other results in Al2O3 doped withdifferent impurities, in which the mechanical behaviour is explained in termsof the change of the ionic bonding strength between Al and O and thecovalent bonding between Al and the surrounding cations, thus affecting thegrain boundary diffusivity.49–51 This explanation has been outlined in SiO2–TZP doped with several kinds of metal oxides.51–53 This change in the bondstrength has also been used to explain the superplasticity of SiC doped withsmall amounts of boron.54 The doped boron segregates at grain boundaries,removing silicon from its site, thence forming bonds in a local environment,similar to that in the B4C structure.55 This fact enhances the deformation bygrain boundary diffusion.


An improvement of the diffusivity and in consequence of the flow stressand elongation to failure is also found when a secondary glassy phase isadded. This behaviour has been reported in barium and borosilicate doped3Y-TZP.7,56 As mentioned above, non-oxide ceramics such as SiC and Si3N4

are fabricated with sintering aids, thence generating a two-phase material,with a hard phase surrounded by a soft secondary glassy phase. This secondaryphase is the one controlling the mechanical behaviour of these materials, asmentioned when referring to the accommodation processes controllingsuperplasticity.

In the case of Si3N4, glass pockets and thin glass film with thickness ofabout 1 mm often remain at grain boundaries.57 Its plasticity is controlled bythe viscosity of the intergranular glassy phase and the solubility of the crystallinephase in the liquid and have been reviewed in several contributions.8,58,59

The solid solution of silicon nitride with some aluminium-based compoundsor mixture form the so called a¢ or b¢-SiAlON compounds which aresuperplastic by the addition of secondary glassy phase; for example, Rosenflanzand Chen60 reported that Li-doped SiAlON deforms 10 times faster thanSi3N4. A revision of the mechanical properties of these compounds can befound in the literature.8,58 The enhancement of superplastic deformation byintergranular glass phase was also applicable to liquid-phase sintered SiC.61,62

As mentioned in the last two paragraphs, to improve superplasticity it isnecessary to reduce the grain size or to enhance the diffusion process controllingit; however, as was recently shown in YTZP, the reduction of grain size mayproduce a reduction of the diffusivity of the species controlling superplasticity.It has been successively shown that the yttrium in YTZP polycrystals segregatesat grain boundaries and this segregation was the possible cause of the thresholdstress (s0) and could explain quantitatively the dependence of this s0 withtemperature and grain size.14

The segregation of yttrium atoms whose electric charge is different fromthat of the parent ions induces a local density of negative charge produced bythe ¢YZr defects accounting for a local electric field which is screened by thegradient of oxygen vacancies between the bulk and the boundaries. When thegrain size of the polycrystal becomes close to the screening length (nanoscalelength), the electric field can influence the diffusional processes and thecreep equation (16.1) will be multiplied by the following factor:63

al

el

= 1

1 + 4 exp– ( )

3 – 1

rdz eV R

kT dDÊ

Ëˆ¯

ÈÎÍ

˘˚

(16.14)

with l the Debye attenuation length, zD the valence of yttrium, V(R) theelectrical potential and er the relative dielectric constant.

A plot of a versus the average grain size from (16.14) is shown in Fig.16.5. From this it can be seen that the effect of the nanostructured character


of the YTZP specimen is more and more pronounced the bigger l is. Thecreep resistance increases up to a factor of 10 for a grain size around 50 nm,when l is equal to 20 nm.

Whereas an ample number of publications have dealt with submicron-sized YTZP, the number of publications decreases when going down tonanometre size, due mainly to the fact that only recently have fully-densifiednanocrystalline YTZP become available. Recently, improved creep resistancehas been reported in 50 nm YTZP deformed at 1200ºC in agreement withprediction with the model.64,65 Few papers have been published 66–68 onnanocrystalline monoclinic ZrO2; however, as mentioned, in monoclinicmaterials there is no segregation at the grain boundaries and the predictionsof the model developed when segregation occurs cannot be tested. At thispoint, it is necessary that more systematic work on well-defined systems beconducted in order to verify the importance of an electric field in the diffusionprocess when the grain size decreases down to the Debye length scale.

16.5 Applications of superplasticity

The increasing applications of advanced ceramics in technical areas includingaerospace, energy, electronics, biology, etc., often require complex shapes tobe manufactured at low prices. The extensive potential applications, togetherwith the possibility of processing dense ceramics and forming complexstructures by superplasticity, have been the driving force for the appearanceand fast development of a large number of ceramic systems with superplasticcapabilities. Several industrial processes in the metal and polymer industries

0 40 80 120 160 200Grain size (nm)

l = 5 nml = 10 nml = 15 nml = 20 nm

1

10–1

10–2

a

16.5 Plot of a versus grain size for different l values.


have already made use of these high-ductile ceramics. The processing ofdense ceramics includes sheet forming, blowing, stamping, forging and joining.A good example of superplastic forming of different ceramics can be foundin Figure 1 of reference 3. In this figure, flat 1 mm thick discs of Si3N4, 2Y-TZP:Al2O3, 2Y-TZP:Mullite and 2Y-TZP + 0.3% doped with Mn, Fe, Co,Cu and Zn, were stretched with a 6.5 mm radius punch at temperatures andforming times depending on the ceramics.

Sinter forging is a promising technique because densification and net-shaping are achieved simultaneously.69 This technique avoids cavities andvoids because sinter forging is produced by compression. High strength andhigh fracture toughness of Si3N4 have been achieved by superplastic sinterforging due to the reduction of flaw size and grain alignment.70

As previously mentioned, superplasticity is due to GBS accommodatedby diffusional processes; a novel technique to join ceramics, which takesadvantage of both processes, has been developed. When two ceramics incontact are deformed in the superplastic regime (i.e. as soon as GBS isactivated), the grains of one part interpenetrate those of the other, producinga rapid and perfect junction of both parts in shorter times and at lowertemperatures than those commonly required in other conventional processesfor ceramics joining.71 An example of this and further proof that GBS is themechanism of superplasticity is displayed in Fig. 16.6. This figure is composedof two parts: Fig. 16.6(a) displays a set of two pieces of 3Y-TZP joined at1400ºC for 15 min; Fig. 16.6(b) displays an analogous set of two 3Y-TZPswith different grain sizes. In the Fig. 16.6(a), the arrows show the interfaceof the two pieces and it is easy to see how the grains of both parts interpenetrateinto each other.

On the other hand, superplasticity is grain size dependent (Eq. (16.1)):materials with coarse grains will be more creep resistant than those withsmaller grains. Based on this behaviour, it was possible to join zirconialayers of different grain sizes to obtain a multilayer composite with cleanand strong interface and heterogeneous mechanical behaviour at room andhigh temperatures depending on the stress application.72 The same behaviourmay be obtained when two layers of different compositions are joined, becausein yttria-stabilized zirconia (YSZ), grain size is a function of the yttria content.73

This fact is displayed in Fig. 16.7. Figure 16.7(a) is a micrograph showingtwo different layers joined together; Fig. 16.7(b) shows a room-temperatureVickers indentation along the interface, where the crack does not propagatealong the interface but inside the brittle material; and Fig. 16.7(c) is a plot ofthe high-temperature plastic deformation at 1400ºC with the compressionaxis parallel and perpendicular to the interface.

Using this technique, a multilayer made of four different layers 0.5 mmthick of 3YTZP with grain sizes of 0.3, 0.5, 0.8 and 1.0 mm, obtained byannealing the as-sintered ceramics (0.3 mm), labelled from ‘a’ to ‘d’, have


been joined together in a sequence ‘abcdabcd’ (Fig. 16.8(a)) at 1400ºC inorder to form a compound with anisotropic mechanical behaviour dependingon the compression axis parallel or perpendicular to the interface.74 Whenthe compression axis is perpendicular to the interface, all the layers aresubmitted to the same stress (thereby an isostress test), and the layer with thesmaller grain size controls superplasticity. When the compression axis isparallel to the interface, all the layers suffer the same strain (thereby anisostrain test) and the layer with the biggest grain size controls the plasticity(Fig. 16.8(b)). Using the creep model for composites with duplex microstructuredeveloped by French et al.,75 it is possible to fabricate a composite with a

(a)

2 mm

(b)

2 mm

16.6 SEM micrographs of junctions of (a) two layers with the samegrain size (0.3 mm), and (b) two layers with different grain sizes(0.3 and 1 mm).


defined creep resistance by controlling the grain size, the width of the layersand the compression axis; in conclusion, with this technique, it is possible toobtain a defined functionally graded material (FGM).

The use of nano-ceramics as interlayers can reduce drastically thetemperature of joining. A good example can be observed in layers of YTZP,which were joined with an interlayer of 20 nm of YTZP at 1150ºC, 200ºClower than the temperature used without the nano-layer.76

16.7 (a) SEM micrograph of two layers of different composition,3 mol% YTZP (3Y) and 6 mol% YTZP (6Y) (b). Optical micrograph ofthe cross-section of the 3Y–6Y junction with a Vickers indentationand the crack pattern. (c) Stress–strain curve of the 3Y–6Y junctiondeformed at 1400∞C with the stress parallel (isostrain) andperpendicular (isostress) to the interface.

(b)

100 mm

3Y

6Y2 mm

6Y

3Y

(a)

0 2 4 6 8 10Strain (%)

Isostrain

Isostress

3Y–6Y

3Y–6Y

40

30

20

10

0

Str

ess

(MP

a)

(c)

T = 1400∞C


16.6 Future trends

In order to outline the future trends in superplasticity in ceramics, first of allit is necessary to give an answer to the following question: why is superplasticityin ceramics so important? The potential use of these materials in more andmore severe applications makes superplasticity in ceramics an importanttool for their processing, as happened with metals at the beginning of the1960s.

16.8 (a) Schema of the composite obtained with four different layerslabelled a to d. Stress-strain curve of the composite deformed at1400∞C in the isostrain and isostress regime.

(a)

Isostress

Isostrain

a

b

c

d

a

b

c

d

Isostrain

Isostress

0 10 20 30Engineering strain (%)

(b)

30

20

10

0

En

gin

eeri

ng

str

ess

(MP

a)


16.6.1 What properties should be controlled insuperplasticity?

The nature of the grain boundaries when segregation or glassyphases exist

It is well known that the impurities can segregate or precipitate at the grainboundaries; however, it is not known what role they play in superplasticity.For instance, what is the importance of the ionic radius of the impurities, thecharge effect or the binding energy between the host atoms and the impuritiesin the superplastic behaviour? As indicated in the text, the charge effect hasbeen used to justify the threshold stress in YTZP and the binding energy hasbeen used to justify the different behaviour of monolithic ceramics likeYTZP or alumina when doped with different impurities, although the ionicradius has also been used for the same explanation by the same authors.

Type of defects, if any, created during GBS

For instance, dislocations have been shown to play a key role in theaccommodation process in YTZP, justifying the threshold stress in YTZP, incontrast with the hypothesis that this threshold stress is due to the electricfield created by impurity segregation. However, dislocations are notsystematically observed in YTZP; furthermore it was shown that in yttria-stabilized tetragonal zirconia single crystals, the stress necessary to activatedislocations at 1400ºC was over 400 MPa, one order of magnitude higherthan the stresses used during superplastic deformation of YTZP at the sametemperature. It will be necessary to conduct a systematic study of themicrostructure of the monolithic ceramics such as YTZP before and afterdeformation and to correlate their relationship with the superplasticfeatures.

Grain boundary sliding, electron density and binding energy

It is necessary to advance in the knowledge of the superplastic equations, atleast for pure monolithic ceramics and composites with glassy phases, topredict the behaviour for a given material. At this point, probably first-principles simulations of grain boundary sliding and first-principles calculationof electron density and the binding energy between the guest and host atomsin grain boundaries will be of great help in this task.77,78 Interatomic potentialsusing the Embedded Atom Method in conjunction with molecular static anddynamics calculations have been used to study the sliding and migration of(110) symmetric tilt grain boundaries in aluminium;79 although in ceramicsthe grain boundaries are more complicated, it could be interesting to attemptthe same task in order to get better insights about GBS.


16.6.2 Future trends in superplastic ceramics

In the immediate future, the main objective in ceramic superplasticity will bethe search of the right conditions to achieve ‘high strain rate superplasticity’(HSRS) ( ( 10 s ).–2 –1e ≥ Although this phenomenon has been found in severalceramic compounds and several inputs have been outlined to achieve it, weare still far from knowing what to do to obtain this effect systematically. ThisHSRS will enlarge the applications for ceramics.

On the other hand, improvements in ceramic powder processing technology,the routine preparation of high-purity ceramics of nanometre scale, and thenew techniques for the processing of these powders such as HIP, SPS,microwave furnace, etc., will be the driving forces for a very active study onnanoceramics in the near future, probably opening up new phenomena andnew applications.

Finally, the forthcoming comprehension of superplasticity will demand awell-settled justification of the basic equations for this phenomenon. In orderto achieve this goal, first-principles calculations should be conducted. This isa very challenging task, because the mechanical behaviour of grain boundariesrequires an understanding of the physics involved at many different scales.At this point, simulations at microscopic as well as mesoscopic levels canbecome a useful tool.


The financial support awarded by the Spanish Ministerio de Educación yCiencia through the Project CICYT no. MAT2003-04199-C02-02 isacknowledged.

16.8 References

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14. Gomez-Garcia, D., Lorenzo-Martin, C., Muñoz, A., and Domínguez-Rodriguez, A.,‘Correlation between yttrium segregation at the grain boundaries and the thresholdstress for plasticity in yttria-stabilized tetragonal zirconia polycrystals’, Phil. Mag,2003, 83, 93–108.

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21. Ishihara, S., et al., ‘Stereographic analysis of grain boundary sliding in superplasticdeformation of alumina–zirconia two phase ceramics’, Mater. Trans. JIM, 1999, 40,1158–65.

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26. Spingarn, J.R., and Nix, W.D., ‘Diffusional creep and diffusionally accommodatedgrain rearrangement’, Acta Metal., 1978, 26, 1389–98.


27. Langdon, T.G., ‘The significance of grain boundary sliding in creep and superplasticity’,Metal Forum, 1981, 4, 14–23.

28. Ruano, O.A., Wadsworth, J., and Sherby, O.D., ‘Deformation of fine-grained aluminaby grain boundary sliding accommodated by slip’, Acta Mater. 2003, 51, 3617–34.

29. Jimenez-Melendo, M., and Dominguez-Rodriguez, A., ‘Like-metal superplasticityof fine-grained Y2O3-stabilized zirconia ceramics’, Phil. Mag, 1999, A 79, 1591–608.

30. Langdon, T.G., ‘A unified approach to grain boundary sliding in creep andsuperplasticity’, Acta. Metall. Mater., 1994, 42, 2437–43.

31. Wang, J., Taleff E.M.. and Kovar, D., ‘High-temperature deformation of Al2O3/Y-TZP particulate composites’, Acta Mater., 2003, 51, 3571–83.

32. Raj, R., and Chyung, C.K., ‘Solution-precipitation creep in glass ceramics’, ActaMetal., 1981, 29, 159–66.

33. Wakai, F., ‘Step model of solution-precipitation creep’, Acta Mater., 1994, 42, 1163–72.

34. Melendez-Martinez, J.J., Gómez-García, D., and Domínguez-Rodríguez, A., ‘A criticalanalysis and a recent improvement of the two-dimensional model for solution-precipitation creep: application to silicon nitride ceramics’, Phil. Mag., 2004, 84,2305–16.

35. Chen, I.W., and Hwang, S.L., ‘Shear thickening creep in superplastic silicon nitride’,J. Am. Ceram. Soc., 1992, 75, 1073–9.

36. Kim, B.N., Hiraga, K., Morita, K., and Sakka, Y., ‘A high-strain-rate superplasticceramics’, Nature, 2001, 413, 288–91.

37. Yonn, C.K., and Chen, I.W., ‘Superplastic flow of two-phase ceramics containingrigid inclusions–zirconia/mullite composites’, J. Am. Ceram. Soc., 1990, 73, 1555–65.

38. Mayo, M.J., ‘Processing of nanocrystalline ceramics from ultrafine particles’, Int.Mater. Rev., 1996, 41, 85–115.

39. Shen, Z., Peng, H., and Nygren, M., ‘Formidable increase in the superplasticity ofceramics in the presence of an electric field’, Adv. Mater., 2003, 15, 1006–9.

40. Wilson, J., and Kunz, S.M., ‘Microwave sintering of partially stabilized zirconia’, J.Am. Ceram. Soc., 1988, 71, C40-1.

41. Mizuno, M., Obata, S., Takayama, S., Ito, S., Kato, N., Hirai, T., and Sato, M.,‘Sintering of alumina by 2.45 GHz microwave heating’, J. Eur. Ceram. Soc., 2004,24, 387–91.

42. Chen, I.W., and Wang, X.H., ‘Sintering dense nanocrystalline ceramics withoutfinal-stage grain growth’, Nature, 2000, 404, 168–71.

43. Mimurada, J., Nakano, M., Sasaki, K., Ykuhara, Y., and Sakuma, T., ‘Effect of cationdoping on the superplastic flow in yttria-stabilized tetragonal zirconia polycrystals’,J. Am. Ceram. Soc, 2001, 84, 1817–21.

44. Nakatani, K., Nagayama, H., Yoshida, H., Yamamoto, T., and Sakuma, T., ‘Theeffect of grain boundary segregation on superplastic behavior in cation-doped 3Y-TZP’, Scripta Mater., 2003, 49, 791–5.

45. Morita, K., Hiraga, K., and Kim, B.N., ‘Effect of minor SiO2 addition on the creepbehaviour of superplastic tetragonal ZrO2’, Acta Mater., 2004, 52, 3355–64.

46. Sato, E., Morioka, H., Kuribayashi, K., and Sundararaman, D., ‘Effect of smallamount of alumina doping on superplastic behaviour of tetragonal zirconia’, J. Mater.Sci., 1999, 34, 4511–18.


47. Yoshida, H., Okada, K., Ikuhara, Y., and Sakuma, T., ‘Improvement of high-temperaturecreep resistence in fine-grained Al2O3 by Zr4+ segregation in grain boundaries’, Phil.Mag. Lett., 1997, 76, 9–14.

48. Wakai, F., Nagano, T., and Iga, T., ‘Hardening in creep of alumina by zirconiumsegregation at the grain boundary’, J. Am. Ceram. Soc., 1997, 80, 2361–6.

49. Yoshida, H., Yamamoto, T., Ikuhara, Y., and Sakuma, T., ‘A change in the chemicalbonding strength and high-temperature creep resistance in Al2O3 with lanthanoidoxide doping’; Phil. Mag, 2002, A82, 511–25.

50. Yoshida, H., Ikuhara, Y., and Sakuma, T., ‘Grain boundary electronic structuralrelated to the high-temperature creep resistance in polycrystalline Al2O3’, Acta Mater.,2002, 50, 2955–66.

51. Ikuhara, Y., Yoshida, H., and Sakuma, T., ‘Impurity effects on grain boundary strengthin structural ceramics’, Mater. Sci. Eng., 2001, A319-321, 24–30.

52. Thavorniti, P., Ikuhara, Y., and Sakuma, T., ‘Microstructural characterization ofsuperplastic SiO2-doped TZP with a small amount of oxide addition’, J. Am. Ceram.Soc., 1998, 81, 2927–32.

53. Ikuhara, Y., Yamamoto, T., Kuwabara, A., Yoshida, H., and Sakuma, T., ‘Structureand chemistry of grain boundaries in SiO2-doped TZP’, Sci. Tech. Adv. Mater, 2001,2, 411–24.

54. Shinoda, Y., Nagano, T., Gu, H., and Wakai, F., ‘Superplasticity of silicon carbide’,J. Am. Ceram. Soc., 1999, 82, 2916–18.

55. Gu, H., Shinoda, Y., and Wakai, F., ‘Detection of boron segregation to grain boundariesin silicon carbide by spatially resolved electron energy-loss spectroscopy’, J. Am.Ceram. Soc., 1999, 82, 469–72.

56. Imamura, P.H., Evans, N.D., Sakuma, T., and Mecaryney, M.L., ‘High temperaturetensile deformation of glass-doped 3Y-TZP’, J. Am. Ceram. Soc., 2000, 83, 3095–99.

57. Clarke, D.R., ‘On the equilibrium thickness of intergranular glass phases in ceramicmaterials’, J. Am. Ceram. Soc., 1987, 70, 15–22.

58. Wilkinson., D.S., ‘Creep mechanisms in multiphase ceramic materials’, J. Am. Ceram.Soc., 1998, 81, 275–99.

59. Wakai, F., Kondo, N., and Shinoda, Y., ‘Ceramics superplasticity’, Curr. Opin. Sol.Sta. Mater. Sci., 1999, 4, 461–5.

60. Rosenflanz, A., and Chen, I.W., ‘Classical superplasticity of SiAlON ceramics’, J.Am. Ceram. Soc., 1998, 81, 713–16.

61. Wang, C.M., Mitomo, M., and Emoto, H., ‘Microstructure of liquid phase sinteredsuperplastic silicon carbide ceramics’, J. Mater. Res. 1997, 12, 3266–70.

62. Nagano, T., Gu, H., Shinoda, Y., Zhan, D., Mitomo, M., and Wakai, F., ‘Tensileductility of liquid-phase sintered b-silicon carbide at elevated temperatures’, Mater.Sci. Forum, 1999, 304–6, 507–12.

63. Gomez-Garcia, D., Lorenzo-Martin, C., Muñoz, A., and Domínguez-Rodriguez, A.,‘Model of high-temperature plastic deformation of nanocrystalline materials:Application to yttria tetragonal zirconia’, Phys. Rev., 2003, B67, 144101–1–8.

64. Gutierrez-Mora, F., Jimenez-Melendo, M., Domínguez-Rodriguez, A., and Chaim,R., ‘High temperature mechanical behavior of YSZ nanocrystals’, Key Eng. Mater,2000, 171–4, 787–92.

65. Gutierrez-Mora, F., Domínguez-Rodriguez, A., and Jimenez-Melendo, M., Chaim,R., and Hefetz., M., ‘Creep of nanocrystalline Y-SZP ceramics’, NanostructuredMater., 1999, 11, 531–7.


66. Roddy, M.J., Cannon, W.R., Skandan, G., and Hahn, H., ‘Creep behaviour ofnanocrystalline monoclinic ZrO2’, J. Eur. Ceram. Soc. 2002, 22, 2657–62.

67. Yoshida, M., Shinoda, Y., Akatsu, T., and Wakai, F., ‘Deformation of monoclinicZrO2 polycrystals and Y2O3-stabilized tetragonal ZrO2 polycrystals below themonoclinic–tetragonal transition temperature’, J. Am. Ceram. Soc., 2002, 85, 2834–6.

68. Yoshida, M., Shinoda, Y., Akatsu, T., and Wakai, F., ‘Superplasticity-like deformationof nanocrystalline monoclinic zirconia at elevated temperatures’, J. Am. Ceram.Soc., 2004, 87, 1122–5.

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Part V

Non-oxide ceramic composites

459

461

17.1 Introduction

An interface is a meeting surface between two dissimilarly oriented perfectcrystals or between two chemically different crystals, i.e., a crystallographicor chemical discontinuity. In terms of crystal structure and chemistry ofadjacent crystals, interfaces can be classified into homophase and heterophaseboundaries (Cahn and Kalonji, 1982; Finnis and Rühle, 1993). Homophaseboundaries form between grains of identical crystal structure and composition,e.g., grain boundaries, twin boundaries, domain boundaries and stackingfaults. Heterophase boundaries form between regions of different crystalstructure and/or chemical composition, e.g., interfaces between co-existingpolytypes, between metal-ceramic joints and most interfaces between thematrix and the reinforcing agents in composite materials. Hence, the term‘interface’ is a general term, whereas the terms ‘grain boundary’ and ‘interphaseboundary’ are more specialised terms used to distinguish between the twotypes of interface.

In contrast to grain boundaries and interphase boundaries in metals, inwhich grains make intimate contact at the atomic level regardless of theorientation relationship across the boundary (Sutton and Balluffi, 1995),grain boundaries in non-oxide engineering ceramics and interphase boundariesin composites invariably contain thin films of extraneous material whicharise through the presence of impurities and additives present in the startingpowders used to make the ceramics and composites (Kleebe et al., 1992;Turan and Knowles, 1995). Here, we will use the term ‘intergranular film’ todescribe a phase at an interface. In the context of the non-oxide particulateceramic composites that we will discuss here, such films are typically1–2 nm thick. This range of thickness contrasts with the thickness of chemicallydistinct regions between the matrix and the reinforcing fibres in fibre-reinforcedceramics where such regions can have thicknesses between 50 nm and a fewmicrometres (see, for example, Kumar and Knowles, 1996a, 1996b). If thisintergranular film is amorphous, a distinction has to be made between

17Interfaces in non-oxide ceramic

composites

S T U R A N, Anadolu University, Turkey andK M K N O W L E S, University of Cambridge, UK


segregation of chemical species and a genuine intergranular film. Rühle etal. (1984) have suggested that an intergranular film forms when the densityof any segregated impurities is larger than one atomic layer, otherwise theterm segregation is more appropriate. Terms such as ‘glassy films’ to describeintergranular films should be used with caution, since there is some evidencethat these films are more ordered than glass at triple junctions (Marion et al.,1987).

Intergranular films have traditionally been viewed as a problem for themechanical properties of engineering ceramics and composites. Typicallythese films have a lower fracture toughness than the surrounding grains, witha correspondingly lower fracture stress. In addition, they tend to have lowersoftening points or melting points than the more refractory matrix grains.The more amorphous phase there is present in an engineering ceramic or anengineering composite, the more it will be expected to creep at hightemperatures, because of the exponential decrease in the viscosity of amorphousphase with increasing temperature. For example, Choi et al. (2004) foundthat silicon carbide (SiC) ceramics are able to retain their strength at 1500∞Cwhen AlN and either Sc2O3 or Lu2O3 are present in the starting powdersbecause the grain boundaries which form are free from intergranular films.However, intergranular films can be beneficial. At temperatures below 1000∞C,Pezzotti (1993), Becher et al. (1998) and Sun et al. (1998) have all reportedthat intergranular films can be used very effectively to enhance the matrixfracture toughness by controlling the interface chemistry. For example, it hasbeen shown that the fracture toughness of in-situ reinforced silicon nitride(Si3N4) ceramics can be optimised by a suitable choice of sintering aids tocontrol the composition of the grain boundary films as well as the surfacechemistry of the adjoining grains (Becher et al., 1998; Sun et al., 1998). Ifthe grain boundaries between adjacent matrix grains in in-situ reinforcedsilicon nitride are too strong, cracks will propagate through the matrix siliconnitride grains rather than be deflected along the more tortuous grain boundaries.Such a silicon nitride engineering ceramic will not have an enhancement ofits fracture toughness over more conventional dense silicon nitride withequiaxed grains. However, if the grain boundaries between adjacent matrixgrains are too weak, the silicon nitride will have a low strength, even if it hasa higher fracture toughness than a more conventionally made silicon nitride.Thus, interfacial structure and interfacial chemistry both play important rolesin the toughening and strengthening response of the material. Amorphousintergranular films also play an important role in both the b (or 3C) Æ aphase transformation in SiC (Moberlychan et al., 1998) and the reverse a Æb (or 3C) phase transformation (Turan and Knowles, 1996a).

In this overview, we will first discuss how transmission electron microscopy(TEM) techniques can be used to determine the presence or absence ofintergranular amorphous phases at interphase boundaries in structural

Interfaces in non-oxide ceramic composites 463

engineering ceramics and non-oxide composites. We will show how theobserved distribution of thicknesses of the order of 1-2 nm for such filmscan be understood in terms of theoretical models of the attainment of anequilibrium film thickness from suitable competing attractive and repulsiveforces at interphase boundaries. We will then discuss the evidence for andagainst the development of preferred orientation relationships and good latticematching at intergranular and interphase boundaries, because suchconsiderations are also relevant for the development of models for the strengthand toughness of engineering ceramics and composites. Finally, we willexamine future trends within this research area.

17.2 Assessment of the accuracy of TEM techniques

for the detection and measurement of film

thickness at interfaces

Detailed characterisation of interfaces requires high magnifications and highinstrumental resolutions. Usually, this precludes the use of techniques suchas scanning electron microscopy, requiring instead more specialist TEM orscanning transmission electron microscopy (STEM) techniques. TEMtechniques used to date for the characterisation of interfaces in non-oxideceramic composites are bright field imaging, dark field imaging, diffusedark field imaging from the part of reciprocal space into which there isscattering from the intergranular amorphous phase, weak beam dark fieldimaging, diffraction pattern analysis, Fresnel defocus imaging and highresolution transmission electron microscopy (HRTEM) (Clarke, 1979; Cinibulket al., 1993a, 1993b; Turan, 1995; Kleebe, 1997, 2002). A recent elegantdemonstration of the capabilities of STEM has been shown by Shibata et al.(2004) who used aberration corrected Z-contrast STEM to examine thesegregation of lanthanum atoms to intergranular films at grain boundaries insilicon nitride.

For these techniques, the interface under examination is oriented edge-onso that it is parallel to the electron beam. For HRTEM, grains at either sideof the interface must be diffracting strongly, ideally with the electron beamparallel to a low-index zone of each grain, whereas in Fresnel fringe analysisthe two grains must be diffracting weakly. Diffuse dark field imaging, HRTEMand Fresnel defocus imaging are particularly useful for characterising interfacescontaining intergranular films, but, in addition, analytical electron microscopytechniques such as energy dispersive X-ray (EDX) and parallel electronenergy loss spectroscopy (PEELS) need to be used to obtain chemicalinformation about the intergranular films.

There is no single experimental technique capable of characterising aninterface structure fully. Therefore, it is desirable to use techniques whichare complementary in order to maximise the available information as well as


the reliability of the results. A comprehensive study by Cinibulk et al. (1993a)in which diffuse dark field imaging, HRTEM and Fresnel defocus imagingwere all compared concluded that HRTEM gives an accuracy of ±1 Å forintergranular film thickness measurements if grains either side of the interphaseboundary have at least a one-dimensional lattice image to distinguish themfrom the intergranular amorphous phase. By comparison, the Fresnel techniquepredicts a thickness 20-35% higher than HRTEM, and diffuse dark fieldmeasurements predict thicknesses some 50-100% higher than those measuredby HRTEM. Similar trends were first noted by Clarke (1979).

An example of the use of Fresnel defocus imaging, diffuse dark fieldimaging and HRTEM to examine the same SiC–SiC interface is shown inFig. 17.1 (Turan, 1995). All three techniques clearly indicate the presence of

3C SiC

3C SiC

a-SiC

(a)

(b)

(c)

21 Å

20 nm

17.1 (a) Underfocus, (b) overfocus Fresnel fringe images obtainedfrom an edge-on interface between SiC grains, (c) a diffuse dark fieldimage, (d) an HRTEM image from the same interface, and (e) a plotof Fresnel fringe spacing against defocus for a part of the interfacewhere two 3C SiC grains are separated with an intergranular filmshown in (a) and (b). The contamination and damage evident in(a)–(d) arises from the time required for the edge-on alignment of theinterface and also because of the high voltages used in HRTEM.


an intergranular amorphous film at this interface. The images shown in Fig.17.1 were taken on two different transmission electron microscopes, sincethe microscope on which the HRTEM work was undertaken did not have asmall enough objective aperture to obtain high-contrast Fresnel and diffusedark field images. Using the methodology described by Cinibulk et al. (1993a),the thickness of the intergranular film was measured to be approximately13 ± 1 Å from HRTEM images well away from the small triple junctionbetween two 3C grains and a-SiC, but it will be apparent from Fig. 17.1(d)that the thickness varies as a function of position along the interface, particularlyin the vicinity of the triple junction.

10 nm

3C SiC

3C SiC

a-SiC

(d)

5.5

4.5

3.5

2.5

1.5

0.5

Frin

ge

spac

ing

(n

m)

–6000 –4000 –2000 0 2000 4000 6000Defocus (nm)

(e)

17.1 Continued


One of the advantages of HRTEM over the other techniques in measuringthe intergranular film thickness accurately is that an internal calibration ispossible for every single image, whether originally in the form of a negativeor a digitised image. Other, more general, calibration methods are usuallyneeded for the two other techniques. However, if phases are present with avery large interplanar spacing which can be suitably oriented, such as the(0001) planes of 6H polytype of a-SiC which have an interplanar spacing of~15 Å, as in Fig. 17.1(d), lattice fringes from these (0001) planes will bereadily visible in any of the recent generation of transmission electronmicroscopes at the typical magnifications of 100 000–150000 used for Fresnelimaging. Thus, an image at the magnification used for Fresnel imagingconditions with these lattice fringes present can be used to calibrate themagnifications used during subsequent Fresnel imaging of the interfaces andduring the eventual image processing.

The through-focal series for Fresnel imaging analysis for the interface inFig. 17.1 contained 21 negatives in total taken from underfocus (DF = -3830nm) to overfocus (DF = +3830 nm). Six of these images were used to find theexact position of the Gaussian focus, where DF = 0 nm, through an examinationof the intensity of the Fresnel fringes at the interface. Unfortunately, lowfringe visibility around Gaussian focus prevents accurate measurements ofthe Fresnel fringe spacings at low defocus values. The remaining 15 fringespacings were then measured and plotted against their corresponding defocusvalues (Fig. 17.1(e)).

The film thickness obtained from Fig. 17.1(e) by extrapolating the highdefocus values to low defocus values so that the underfocus and overfocuscurves meet corresponds to a film thickness of ~16 Å. This is ~23% higherthan the value obtained from HRTEM measurements.

The film thickness was determined to be ~21 Å from the diffuse dark fieldimage which is ~ 60% larger from that of HRTEM measurements and is alsoconsistent with the results obtained by Clarke (1979) and Cinibulk et al.(1993a). Diffuse dark field imaging is also the technique most likely toproduce artefacts, most notably from preferential etching from ion beamthinning of the interfaces and subsequent sputter deposition or damage inthese regions.

The question of whether modern-day SEMs can be used to infer thepresence or absence of intergranular films at grain boundaries has recentlybeen addressed by Kleebe (2002). Kleebe characterised two different non-oxide ceramics, Si3N4 and SiC, with respect to their grain boundary structureusing both SEM and TEM. SEM imaging of plasma etched surfaces revealeda characteristic bright contrast along the interfaces for both ceramics, suggestingthe presence of an intergranular amorphous film. HRTEM studies of theSi3N4 sample confirmed that these fine bright lines along grain boundariesrepresented intergranular amorphous films separating Si3N4 matrix grains.


However, when HRTEM was employed on the SiC samples, which showeda similar contrast variation across SiC grain boundaries in the SEM, thepresence of residual intergranular films was not detected even at the triplejunctions. Hence, Kleebe concluded that SEM imaging and Fresnel fringeTEM imaging alone do not enable a safe conclusion to be drawn aboutinterface wetting in ceramic polycrystals.

A further aspect of interfaces in engineering ceramics was addressed byGu and Shinoda (2000). These workers used spatially resolved electron energy-loss spectroscopy (EELS) analysis and spatially resolved energy-loss near-edge structures (ELNES) analysis to characterise interfaces in a hot-isostatically-pressed SiC material whose powders contained ~3 wt% freecarbon and 1 wt% amorphous boron. They demonstrated that the structuralwidth and the chemical width of general boundaries were not the same. TheirHRTEM observations did not detect the existence of any amorphous film atsuch grain boundaries. Instead, each grain boundary had a core structure1–2 atomic planes wide in which B–C and Si–O bonds were detected byEELS and ELNES, as well as Si–C bonds subtly different from the Si–Cbonds in the grain interiors. The chemical width of the grain boundaries thatthey inferred from EELS analysis of elemental profiles was visibly widerthan this core region. ELNES analysis distinguished a third extended grainboundary width larger than the chemical width within which the Si–C bondingwas modified from that of the grain interiors, being most strongly modifiedat the grain boundary and decreasing in modification with increasing distancefrom the grain boundary. Thus, in these materials, they concluded thatintergranular films were not seen, but rather segregation of boron and oxygento the grain boundaries, with modification of the Si–C bonding extendingseveral lattice planes into each adjacent SiC grain.

17.3 Wetting, non-wetting and dewetting

behaviour of interphase boundaries in

non-oxide ceramic composites

TEM studies using Fresnel, diffuse dark field, HRTEM, Z-contrast imagingand spatially resolved electron energy-loss spectroscopy techniques haverevealed that, while most of the interfaces in non-oxide engineering ceramicsand composites are covered with intergranular films (Clarke, 1979; Cinibulket al., 1993a, 1993b; Turan, 1995; Gu et al., 1995; Turan and Knowles, 1995;Kleebe, 2002), some interfaces are clearly not (Schmid and Rühle, 1984;Turan, 1995; Knowles and Turan, 1996; Turan and Knowles, 1997, 1999).Interfaces without intergranular films will be considered in more detail inSection 17.6. Examples of wetting and non-wetting behaviour of grainboundaries are shown in Fig. 17.2.

The converse behaviour to wetting, dewetting of a liquid from an interface


to the sample surface, is also occasionally reported in liquid-phase sinteredceramics. For example, such behaviour has been observed on aluminiumnitride (AlN) ceramics when annealed at high temperature in a highly reducingatmosphere (Ueno and Horiguchi, 1989) and in SiAlON ceramics after vacuumheat treatment (Mandal and Thompson, 2000). Kleebe and Pezzotti (2002)characterised the grain-boundary structure of a model SiAlON polycrystalwith nominal composition Si5AlON7 by TEM both in an equilibrium (as-

17.2 HRTEM examples of non-wetting behaviour at interfaces in(a) SiC and (b) Si3N4, and wetting behaviour at an interface in (c)Si3N4. (a) and (c) reprinted from Mat. Sci. Forum, Turan S andKnowles KM, ‘Wetting and non-wetting behaviour of SiC grainboundaries’, 294–296, 313–316 (1999) with kind permission of Trans.Tech. Publications.

(a) (b)

Si3N4

3C SiC

5 nm 2 nm

(c)

10 nm

Si3N4

Si3N4

Si3N4

Si3N4

a-SiC


processed) state at room temperature and after quenching from elevatedtemperature. They found that in the equilibrated low-temperature microstructureamorphous phases existed only at multigrain junctions, and not at grainboundaries. However, samples of this model polycrystal heated to 1380∞Cand rapidly quenched showed wetted grain boundaries.

17.4 Equilibrium film thickness at interphase

boundaries

Other than the wetting, non-wetting and dewetting behaviour of interfaces,the other striking TEM observation is that the thickness of intergranularfilms seems to be relatively constant from one interface to another in anygiven sample, but that it varies from material to material.

Theoretical considerations by Clarke and co-workers (Clarke, 1987; Clarkeet al., 1993) show that an equilibrium film thickness arises from the competitionbetween attractive dispersion forces determined by the dielectric propertiesof the grains and repulsive disjoining forces which can be steric forces and/or double-layer forces. Wetting will occur when the solid–solid boundaryenergy, gb, is less than that of the wetted boundary, 2gl, where gl is the liquid–solid interfacial energy (Clarke, 1985), provided that there is a suitable sourceof liquid, for example as a consequence of liquid-phase sintering at hightemperatures.

The most general equilibrium condition which applies to a thin filmsandwiched between two phases when there is an applied pressure, P, and acapillary pressure, PCAP, is

P + PCAP + PDISP + PST + PEDL + PADS + PHB = 0 (17.1)

where PDISP is an attractive dispersion force per unit area of interface arisingfrom van der Waals forces, PST is a repulsive steric force per unit area, PEDL

is a repulsive electrical double-layer force per unit area, PADS is a repulsiveforce per unit area arising from the effects of any solute absorption, and PHB

is an attractive force per unit area arising from any hydrogen bonding present(Clarke, 1987; Clarke et al., 1993).

Clarke (1987) examined the form of Eq. (17.1) for the situation wherePEDL = PADS = PHB = 0, in which case the repulsive force per unit areaenabling an equilibrium film thickness to arise is simply PST. Subsequently,Clarke et al. (1993) examined the situation where PEDL π 0, PADS = PHB =0 for zero and finite values of PST, concluding that it is only under certainrestricted conditions that it is plausible for an electrical double-layer force tocontribute significantly to the total repulsive force.

If the only significant repulsive force per unit area is PST, as is likely tobe the case for interphase boundaries between non-oxide engineering ceramicssuch as SiC, Si3N4 and h-BN, the net force, F, per unit area pushing twograins together becomes


F P P P P AL

aL

= + + + = + + 6

– sinh ( /2 )CAP DISP ST CAP 3

2

2P Pp

hx

(17.2)

where L is the thickness of the film, A is the Hamaker constant determiningthe magnitude of the attractive dispersion force, x is a molecular correlationdistance and ah2 is a constant in the term representing the repulsive stericforce per unit area which is the free energy difference between ordered anddisordered states of the film (Clarke, 1987). If L is the equilibrium thickness,F is simply zero.

The trend which arises from a consideration of Eq. (17.2) is that the lowerthe value of the Hamaker constant, A, the higher the equilibrium film thickness(Clarke, 1987). Striking confirmation experimentally of such a trend hascome from work in which the local Hamaker constants in silicon nitrideceramics have been determined from spatially resolved-valence electron energy-loss spectroscopy (French et al., 1998). Conversely, if A is too large, thenthere will be no thickness L for which Eq. (17.2) is satisfied.

For a suitably high critical value of A, this theoretical model predicts alower limit on the equilibrium thickness that can be observed. This lowerlimit on L, Lmin, is defined by the conditions F = 0 and dF/dL = 0 since fora stable film F = 0 and dF/dL ≥ 0 (Clarke, 1987). Various solutions to theseconditions have been examined by Knowles and Turan (2000). In the absenceof capillary pressure and external pressure, Lmin = 2.58x. Using reasonableestimates for x, Knowles and Turan estimate Lmin to be ≥6.50 Å. That inpractice the observed intergranular film thicknesses are typically of the orderof 1–2 nm in non-oxide engineering ceramics indicates that the relevantHamaker constants for ceramics are significantly lower than the critical value.

Israelachvili (1992) discusses the various approximate and analyticformalisms for A, with the conclusion that where two macroscopic isotropicphases 1 and 2 interact across an isotropic medium 3, a suitable approximationto the relevant Hamaker constant, A132, valid for L greater than moleculardimensions is:

A k TB1321 3 2 3

1 3 2 3 = 3

4( – )( – )( + )( + )

e e e ee e e e

+ 3

8 2

( – )( – )

( + ) ( + ) ( + ) + ( + )12

32

22

32

12

32 1/2

22

32 1/2

12

32 1/2

22

32 1/2

h n n n n

n n n n n n n nen

[ ](17.3)

where kB is Boltzmann’s constant, T is absolute temperature, e1, e2 and e3 arethe static dielectric constants of the three phases, ne is the characteristicabsorption frequency taken to be the same for all three materials, h is Planck’s


constant and n1, n2 and n3 are the refractive indices of the three materials,extrapolated to zero energy, or equivalently, zero frequency. Of the two

terms on the right-hand side of Eq. (17.3), the first term can be at most 34

kBT. For most situations of practical interest, this term is significantly smallerthan the second term.

It is usual in calculations using Eq. (17.3) for the common characteristicabsorption frequency ne to be assigned a value of 3 ¥ 1015 s–1 (see, forexample, Israelachvili, 1992; French, 2000). Horn and Israelachvili (1981)have derived a slightly more complex form of Eq. (17.3) for the situationwhere materials 1 and 2 have the same absorption frequency but material 3has a different value of absorption frequency, and Prieve and Russel (1988)have derived a form of A132 for the most general situation where the threematerials have different absorption frequencies n1, n2 and n3.

Knowles and Turan (2000) and Knowles (2005) have used the approachof Parsegian and Weiss (1972) to examine the effect of anisotropy on Hamakerconstants. Such calculations are relevant for materials such as h-BN andrutile, TiO2, which exhibit strong anisotropy in their refractive indices becauseof their crystal structure. They make little difference to predicted values ofHamaker constants as a function of grain orientation for materials such asSi3N4 and SiC which, by comparison, exhibit modest values of birefringence.

The analysis of Knowles and Turan (2000) of h-BN–amorphous silica–3CSiC interfaces showed that Eq. (17.3) could be used to calculate values of theHamaker constant as a function of the orientation of h-BN with respect to aplanar interface containing a thin amorphous silica film, provided that theeffective values of static dielectric constant and refractive index for h-BN,eh–BN and nh–BN respectively, were taken to be

e e e e e qh–BN o

1/2e o e

2 1/2

= ( ) + ( – )sin

2ÏÌÓ

¸˝˛

(17.4)

and

n n n n nh–BN2

o2

e2

o2

e2

2 = 1

2 + 1

2 + ( – ) sin

4q (17.5)

where q is the angle between the normal to the interface plane and the normalto the (0001) h-BN planes, no is the refractive index of the ordinary rays in h-BN, and ne is the refractive index of the extraordinary ray. Forh-BN no >> ne and eo >> ee. However, the dominance of the second term on theright in Eq. (17.3) means that A is most sensitive to the difference between theeffective refractive index, nh-BN, of h-BN and the refractive index of amorphoussilica. Since nh-BN is always greater than the refractive index of amorphoussilica for all values of q, it follows that for h-BN–amorphous silica–3C SiCinterfaces A is predicted to increase as q increases from 0∞ to 90∞. Detailed


calculations by Knowles and Turan showed that this increase was significant:at q = 0∞, A was calculated to be 113 zJ, increasing to 139 zJ at q = 90∞. Thelimited experimental observations they were able to make on h-BN–3C SiCinterfaces containing amorphous silica-rich intergranular films were consistentwith this trend in A, i.e., those interfaces parallel to (0001) h-BN planes hadthicker intergranular films of 12 ± 1 Å, whereas the intergranular film at aninterface where the (0001) h-BN planes made an angle of 68∞ with respect tothe interface plane was noticeably smaller, ~ 8.5 Å (Fig. 17.3).

(a)

(b)

(c)

h-BN

12 Å

3C SiC

h-BN

12 Å

3C SiC

h-BN

8.5 Å

3C SiC

17.3 HRTEM observations of three differently misoriented interphaseboundaries between hexagonal boron nitride (h-BN) and 3C siliconcarbide (3C SiC) grains showing an orientation dependence onequilibrium film thickness. In (a) and (b) the (0001) of the highlyanisotropic h-BN are parallel to the interface, whereas in (c) theymake an angle of 68∞ with the interphase boundary (reprinted fromUltramicroscopy, Knowles KM and Turan S, ‘The dependence ofequilibrium film thickness on grain orientation at interphaseboundaries in ceramic–ceramic composites’, 83(3/4) 245–259 (2000)with kind permission of Elsevier Science).


17.5 Effect of intergranular film composition on

equilibrium film thickness

As we have already noted in Section 17.4, an examination of the behaviourof Eq. (17.2) shows that the equilibrium film thickness, L, increases withdecreasing Hamaker constant. Furthermore, the dominant term in Eq. (17.3)for the Hamaker constant for two isotropic media interacting across a thirdisotropic medium is the one dependent on the refractive indices of the threemedia, n1, n2 and n3 respectively. Thus, if n1 and n2 are both greater than n3,as will be the case for silica-rich films between grains of either SiC or Si3N4,it is obvious from Eq. (17.3) that the higher the refractive index of theintermediate third medium, the lower the Hamaker constant will be, and, inturn, the higher L will be. Conversely, a decrease in the molecular correlationlength, x, will cause a decrease in L.

The refractive index of the intermediate third medium is determined by itschemical composition. Although the exact chemical compositions ofintergranular films have to be known for accurate Hamaker constantcalculations, experimental limitations of both interference with adjacent grainsand beam broadening during chemical analysis in the electron microscopeshave in practice meant that it has not been possible to ascertain these withconfidence. The chemical analysis that has been undertaken has shown thatmost of the films in ceramics are amorphous silica mixed with a very smallamount of metal ion contamination from the starting powders. In addition,there are some indications that intergranular silica films can, and will, containchemical elements from sintering additives and incomplete chemical reactionsduring processing. For example, nitrogen and yttrium enrichment would beexpected in a predominantly silica-rich intergranular film between Si3N4 andY2Si2O7 grains, Si–C–O glasses rather than amorphous silica would be expectedbetween SiC grains, and Si–C–O–N glasses rather than amorphous silicawould be expected between SiC and Si3N4 grains.

The effect of the nitrogen content on the refractive index of M-SiAlONglasses reported in Lewis (1989) clearly shows that when the nitrogen contentincreases, the refractive index of the glass also increases. It is also believed(Lewis, 1989) that by replacing oxygen, nitrogen increases the density ofSiAlON materials, because nitrogen has a superior ability to oxygen forcross-linking the network structure of oxynitride glass - nitrogen atoms arecapable of bridging three network tetrahedral groups, whereas oxygen atomsare capable of bridging only two. In this context, it is relevant to examine asituation where two Si3N4 grains are separated by a supposedly pure silicafilm and to assume instead that 10 mol% nitrogen is in it, undetected, as hasbeen claimed by Vetrano et al. (1993). From Lewis (1989), the effect of 10mol% nitrogen is predicted to increase the refractive indices of the SiAlONglasses by ~4%. This increase, while modest, increases the refractive index


of silica from 1.45 to 1.51 and decreases the predicted Hamaker constantfrom 82 to 65 zJ. Thus, a small amount of nitrogen enrichment in the filmcan radically affect the Hamaker constant and therefore the equilibrium filmthickness.

Kleebe (1997) has summarised experimental results on equilibrium filmthicknesses from a number of studies on interfaces in Si3N4 ceramics anddiscussed them in the light of the competing attractive and repulsive forcespresent across the films and their chemical composition. One of these studiesin which calcium dopant was deliberately introduced into the processingprocedure is of particular note. In this study, originally reported by Tanakaet al. (1994), the equilibrium film thickness at Si3N4 grain boundaries wasshown to be very sensitive to the addition of small quantities of calcium ions.The film thickness was found to be 10 Å for undoped material and 7, 11 and15 Å for samples doped with 80, 220 and 450 ppm Ca respectively. Asexpected, calcium was detected at the grain boundaries in the doped specimens.Tanaka et al. (1994) were able to account qualitatively for their results byconsidering the effect of calcium ions on the silica network structure andalso the development of a repulsive electrical double-layer force, PEDL,through the incorporation of calcium ions as adsorbed species on the grainsurfaces. The qualitative argument they advocated was that for small additionsof calcium, any repulsive electrical double-layer force is more than compensatedby the disruption of the silica network which lowers the magnitude of themolecular correlation distance, x, so that as a result the thickness of the film,L, decreases relative to calcium-free Si3N4 grain boundaries. For larger additionsof calcium, Tanaka et al. (1994) argue that PEDL will increase, so that L willbegin to increase with calcium concentration.

With the increase in the speed and capacity of modern computers, it isnow possible to undertake atomistic simulations of the structure of intergranularfilms. The recent work of Garofalini and Luo (2003) and Su and Garofalini(2004a, 2004b) on calcium silicate intergranular films in silicon nitride isdirectly relevant to the experimental work of Tanaka et al. (1994). Thus, forexample, Su and Garofalini (2004a) examined film compositions of 1.5 mol%calcium, equivalent to 80 ppm calcium doping in the bulk material. Theirwork suggests that an electrical double layer does not form at this level ofcalcium, as they found no evidence of calcium segregation to the grainsurfaces. However, their simulations do show evidence for ordering withinthe intergranular film induced by the crystal surfaces and the incorporationof nitrogen into the films. The study of Garofalini and Luo (2003) on higherconcentrations of calcium in the intergranular layer also failed to show evidencefor segregation of calcium to the grain surfaces, although it did find evidencefor alignment of the calcium ions within the films parallel to the grain surfaces,which weakened the interface by causing a decrease in the number of Si–O–Si bonds across the interface. Interestingly, none of these recent atomistic


studies attempted to interpret the results of their atomistic simulationsquantitatively in terms of the various terms in the force balance of Eq. (17.1)– this suggests that much still remains to be done to understand how thedetail of the intergranular film chemical composition at an atomic leveldetermines the equilibrium film thickness.

17.6 Crystallography of interphase boundaries

In contrast to the number of research studies concerning the thickness andchemistry of intergranular films in non-oxide ceramics and ceramic composites,there have been relatively few studies on the crystallography of either grainboundaries or interphase boundaries in these materials. On the basis of TEMevidence Niihara et al. (1990) suggested that low energy interfaces developpreferentially in Si3N4 grains containing small SiC precipitates, withoutspecifying uniquely the orientation relationships and interface planes thatthey observed. A more detailed two-part TEM study by Unal et al. (1992)and Unal and Mitchell (1992) on chemically vapour deposited Si3N4 grownon single crystal SiC found two characteristic orientation relationships betweenthe Si3N4, which deposited itself as a-Si3N4, and the substrate SiC, whichnear the surface was twinned 3C SiC. The dominant orientation relationshipthat they reported can be described approximately as [101] 3C SiC || [0001]a-Si3N4 with ( (111)) 3C SiC || ( (1010) ) a-Si3N4. Unal et al. accounted fortheir observations in terms of the need to match SiN4 and SiC4 tetrahedrafavouring ( (111)) 3C SiC || (1010) a-Si3N4 || {111} 3C SiC. They accountedqualitatively for small rotations away from this ideal orientation relationshipin terms of the relief of structural mismatch arising from the large misfitbetween the corresponding crystal planes of the two crystal structures.

This orientation relationship was also reported by Pan et al. (1996a) between3C SiC and b-Si3N4 for the ‘type A’ SiC nanoparticles they observed surroundedby b-Si3N4 matrix. b-Si3N4, which has a hexagonal structure with a = 7.6044Å and c = 2.9075 Å (JCPDS-ICDD No. 41-0360) is very similar in structureto a-Si3N4, which has a trigonal structure with a = 7.7541 Å and c = 5.6217Å (JCPDS-ICDD No. 33-1160) and twice the c lattice repeat. Pan et al. alsoobserved that amorphous intergranular films were only seen in some parts ofthe interfaces for the ‘type A’ nanoparticles. In contrast to these ‘type A’nanoparticles, Pan et al. also found other nanoparticles which they designated‘type B’. These nanoparticles had random orientation relationships with respectto the surrounding b-Si3N4 matrices and evidence of substantial amorphousintergranular phase present at the SiC-b-Si3N4 interfaces.

A more detailed study of the crystallography of SiC-Si3N4 interphaseboundaries was undertaken by Turan and Knowles (2000). The compositesthat they investigated containing either 10 or 20 wt% Si3N4, with the balanceSiC. Composites were prepared by hot isostatic pressing. Two different types


of interphase boundary were found in these composites as a consequence ofthe bimodal size distribution of the b-Si3N4 grains formed in the compositesduring the processing operation. In general, interphase boundaries betweensmall intragranular b-Si3N4 precipitates and surrounding SiC grains werefound to be relatively free of intergranular films, whereas interphase boundariesbetween large b-Si3N4 grains and adjacent SiC grains were invariably coveredwith thin intergranular films. Orientation relationships approximating to [110]3C SiC || [0001] b-Si3N4 and (001) 3C SiC || ( (1010) ) b-Si3N4 were foundto dominate between the 3C SiC grains and the small intragranular b-Si3N4

precipitates, the interfaces of which were either clean or relatively devoid ofamorphous intragranular materials (Fig. 17.4). In marked contrast to this,there was no evidence of any favoured orientation relationship between thelarge b-Si3N4 grains and adjacent 3C SiC grains. This dominant approximateorientation relationship is the same as one of the orientation relationshipsreported by Pan et al. (1996b) between 3C SiC particles and a surroundingb-Si3N4 grain. It is closely related to the orientation relationship reported byUnal et al. (1992), Unal and Mitchell (1992) and Pan et al. (1996a): arelatively small rotation of 5.26∞ about the common [110] 3C SiC || [0001]b-Si3N4 direction is required to bring (111) 3C SiC || (0110) b-Si3N4.

(a) 3C SiC

Si3N4

100 nm

(b)

17.4 (a) A dark field image of an interphase boundary between a SiCand Si3N4 precipitate when the electron beam is parallel to [110], (b)its composite diffraction pattern (c) schematic overlapping diffractionpattern, (d) a low magnification HRTEM image showing a Si3N4precipitate embedded in a 3C SiC grain, (e) its composite diffractionpattern, (f) schematic overlapping diffraction patterns,(g)-(i) HRTEM images from IB1, IB6 interphase boundaries and theregion between IB1 and IB6 in (d) (reprinted from Interface Science,Turan S and Knowles KM, ‘The crystallography of interface bondariesbetween silicon carbide and silicon nitride in silicon nitride–siliconcarbide particulate composites’, 8(2/3) 279–294 (2000) with kindpermission of Springer Science and Business Media).


(d)

3C SiC

IB1

IB6 IB2

IB5

IB4

IB3

Si3N4

20 nm

(e)

(c)

002

002

1 1 1

17.4 Continued


(g)

(h)

(i)

5 nm

5 nm

5 nm

Si3N4

3C SiC

17.4 Continued

(f)

m1

m2

Si3N4

3C SiC

002

(3030) (101 0)

220

(2420)

1 1 1 1 1 1

(1210)

1 11 1 1 1

220

002


At present there is no suitable atomistic modelling algorithm which canbe used to examine structures and energies of interfaces between crystallinephases where one or more phase is covalently bonded. Therefore, Turan andKnowles explored the rationale for ‘special orientation relationships’ arisingwhen there is no evidence for the presence of an intergranular film present atSiC-b-Si3N4 interfaces geometrically using the near-coincidence site latticemodel. In such a purely geometric approach, the details of atomic bonding atthe interfaces are necessarily of secondary importance, even though ultimatelythe adoption of any particular three-dimensional orientation relationship andthe energetics of the interface will be determined by the way in which theatoms bond at the interface. Their computations showed that, relative to allpossible orientation relationships between 3C SiC and b-Si3N4 that could beadopted, the dominant orientation relationships between the 3C SiC grainsand the intragranular b-Si3N4 precipitates have low misfits. While obviouscaution must be exercised in the use of purely geometric criteria for lowinterfacial energies (see, for example, Sutton and Balluffi (1987) who concludethat no geometric criterion for low interfacial energy can be regarded aswholly reliable), it must be regarded as significant that the dominant observedorientation relationships between the 3C SiC grains and the intragranularb-Si3N4 precipitates have low misfits, even though it is not possible to inferon the basis of geometry alone that low-energy interfaces occur when theseorientation relationships are adopted.

Turan and Knowles (1997) and Knowles and Turan (2002) have alsoexplored the crystallography of interfaces between b-Si3N4 and hexagonalboron nitride, h-BN, and interfaces between 3C SiC and h-BN, respectively.In the samples they investigated, sub-micron sized h-BN platelets boundedby (0001) h-BN planes occurred as a contaminant. Most of these h-BNplatelets were aligned with respect to SiC grains so that (111) 3C SiC and(0001) a-SiC planes were parallel to (0001) h-BN planes, with [1120] h-BNparallel to either [110] 3C SiC or [1120] a-SiC as appropriate (Fig. 17.5).Clear evidence of thin 12 ± 1 Å amorphous intergranular films were foundat the SiC-h-BN interfaces when there was sufficient liquid available duringthe high temperature processing operation used to manufacture the samples(Turan and Knowles, 1996b). However, in contrast to interfaces between SiCand Si3N4, attempts to rationalise the occurrence of the dominant orientationrelationships between h-BN and SiC were unsuccessful: unrealistically largemisfits would have to be generated within parallel (111) 3C SiC and (0001)h-BN planes in the absence of any amorphous intergranular phase. InsteadKnowles and Turan argued that the {111} 3C SiC planes act as templatesupon which the h-BN basal plane meshes deposit and grow.

Clear orientation relationships were also seen by Turan and Knowles(1997) for interfaces between h-BN and b-Si3N4 (Fig. 17.6). As for interfacesbetween h-BN and SiC, interphase boundaries were dominated by (0001)


h-BN planes, but, significantly, HRTEM suggested that interfaces betweenh-BN and b-Si3N4 were free of amorphous intergranular films, such as the

(a)

100 nm

h-BN

3C SiC 6H SiC

Amorphous phase

(b)

(c)

3C h-BN

12 ± 1 Å

(d)

6H

12 ± 1 Å

h-BN

17.5 (a) A h-BN particle sandwiched between 3C SiC and 6H SiCgrains in a Si3N4–SiC composite made using as-received powders.(b) The diffraction pattern obtained when both SiC grains and theh-BN grain were inside the selected area aperture, showing that inthe h-BN grain and 6H SiC grain the beam direction is parallel to [1120] and that in 3C SiC the beam direction is parallel to [110].Double arrowed spots are 0001 and 1103 2H SiC reflectionsrespectively in the [1120] 2H SiC zone. Weak 1101 and 1103 h-BNreflections arise at the single arrowed positions. (c) The 3C SiC–h-BNinterphase boundary, and (d) the h-BN–6H SiC interphase boundariesshowing thin amorphous intergranular films at both interfaces (Figs17.5(a), (c) and (d) represented from J. Am. Ceram. Sci., Turan S andKnowles KM, ‘Effect of boron nitride on the phase stability andphase transformations in silicon carbide’, 79(12), 3325–3328 (1996)with kind permission of Blackwells Publishing Ltd).


17.6 (a) An example of HRTEM image of an interphase boundariesbetween h-BN particles and adjacent b-Si3N4 grains, showing that theamorphous phase at the triple junction does not seem to extendalong the interphase boundary, (b) and (c) other examples ofintergranular film free interphase boundaries between h-BN and b-Si3N4 (reprinted from J. Eur. Ceram. Soc., Turan S and Knowles KM,‘Interphase boundaries between hexagonal boron nitride andbetasilicon nitride in silicon nitride–silicon carbide particulatecomposites’, 17(15/16), 1849–1854 (1997) with kind permission ofElsevier Science).

(c)

4 nm

h-BN

Si3N4

(a)

h-BN

Si3N4

4 nm

(b)

Si3N4

h-BN4 nm


example shown in Fig. 17.6(c). Knowles and Turan (2002) showed that thetwo orientation relationships found by Turan and Knowles, (i) [11 2 0] h-BN|| [1213] b-Si3N4 and (0001) h-BN || (1010) b-Si3N4, and (ii) [1120] h-BN|| [0001] b-Si3N4 with (0001) h-BN 3.5–4∞ from (1010) b-Si3N4, could bothbe rationalised in terms of low misfits on the basis of the near-coincident sitelattice approach.

Thus, overall, there is strong evidence that non-random orientationrelationships can be adopted in ceramic composites between dissimilar phases.This is most likely when one phase is formed inside the other during high-temperature processing. If there is evidence of remnant intergranular liquidphase at the interface between the two phases, such as at h-BN- SiC interfaces,there is no reason to expect that the interface is of low energy, but thatanother explanation is more appropriate, such as the templating explanationoffered by Knowles and Turan (2002). However, if the interfaces are free ofintergranular film, or if the intergranular film coverage is relatively incomplete,then the analyses of Turan and Knowles (2000) and Knowles and Turan(2002) show that the observed orientation relationships do correlate withthose expected on the basis of relatively low misfits.

17.7 Future trends

The importance that interfaces have in determining structural and functionalproperties of engineering ceramics and composites and the trend towardsnanostructured materials, and thus materials with large areas of grain boundaryper unit volume, will both mean that interfaces will continue to be the objectof intensive research interest. Ultimately, knowledge of the atomistic andelectronic properties of interfaces in non-oxide engineering ceramics andcomposites should lead to the ability to design and control interfacial structure,leading to improvements in mechanical properties such as creep, strengthand fracture toughness.

Substantial progress has already been made experimentally and throughthe use of continuum models in understanding equilibrium film thicknessesof silica-rich films at grain boundaries in engineering ceramics such as Si3N4

and SiC. Progress is also beginning to be made in atomistic modelling ofsuch films, as the recent work of Garofalini and co-workers (Garofalini andLuo, 2003; Su and Garofalini, 2004a, 2004b) has shown.

Experimentally, the technique of measuring Hamaker constants in situusing spatially resolved-valence electron energy loss (SR-VEEL) spectroscopyin the STEM with a 0.6 nm probe (French et al., 1998) represents an importantnew tool for dispersion force and wetting studies (French, 2000). In addition,local variations in intergranular film chemistry and the dispersion forcesthroughout the microstructure of individual grains can now be determinedusing these methods. Other electron microscope-based techniques which are


being developed, and either are already being used or have the potential forfuture use in the study of interfaces in ceramics, include Z-contrast imaging(Shibata et al., 2004; Ziegler et al., 2004), Fourier filtering to remove thelattice fringes from the image and enhance the visibility of intergranularfilms (Maclaren, 2004), electron diffraction with a convergent probe focusedon the intergranular film to obtain information about the local atomic structureof the film (Doblinger et al., 2004), focus-variation phase-reconstructionmethods in HRTEM to image the atomic structure of interfaces at sub-Ångstrom resolution (Ziegler et al., 2003) and electron tomography to obtainthree-dimensional structures (Midgley and Weyland, 2003; Weyland andMidgley, 2004). In addition, advances in atomic force microscopy now enablethis technique to image surfaces of conductors and insulators in vacuum atatomic resolution (Giessibl 2003). Each of these techniques will be able tomake significant contributions to our understanding of interfaces in engineeringceramics and composites.

Modelling studies on interfaces have already made significant advanceswith the ability to undertake molecular dynamics studies (such as the workof Garofalini et al. referred to above) and first-principles calculations (see,for example, Painter et al., 2004). Progress is now being made in understandingthe degree to which intergranular films are ordered and how cations introducedinto the silica-rich films position themselves within the film. Painter et al.(2004) have shown recently that the origin of grain growth anisotropy insilicon nitride ceramics doped with rare earth elements lies in the sitecompetition between the rare-earth cations and silicon for bonding atb-Si3N4 interfaces and within the intergranular film. We can reasonablysurmise that there will be further advances in our understanding from suchcomputer modelling. An area in which progress is still to be made is that ofexamining structures and absolute energies of interfaces between covalentlybonded crystalline phases, such as between grains of silicon nitride or siliconcarbide devoid of any intergranular film, in sharp contrast to interfaces inmetallic materials and ionic solids. Such progress will help in the understandingof the crystallography of interphase boundaries in engineering ceramics andcomposites discussed in Section 17.6.

Overall, therefore, we can confidently predict that the need for detailedknowledge about the behaviour of interfaces in non-oxide ceramics will leadto advances in a number of areas both experimentally and computationally.The relationship between macroscopic continuum models of the wettingbehaviour of intergranular films and the forces keeping intergranular filmspresent at interfaces, and the atomic-level picture gained from high resolutionelectron microscope studies and computer simulations, could be a particularlyfascinating topic – it is apparent that there is substantial scope for progressin this area.


17.8 Further reading

The references in this section are not meant to be an exhaustive list of further researchpapers and books relevant to the work we have described in this chapter. Instead, they area representative selection of material relevant to the work we have described, but havebeen unable to reference in the text, because of the need for conciseness and brevity.Ahn, C.C., (2004), Transmission Electron Energy Loss Spectrometry in Materials Science

and the EELS Atlas, New York, John Wiley & Sons.Angelescu, D.E., Harrison, C.K., Trawick, M.L., Chaikin, P.M., Register, R.A. and Adamson,

D.H., (2004), ‘Orientation imaging microscopy in two-dimensional crystals viaundersampled microscopy’, Appl. Phys. A – Materials Science and Processing, 78 (3):387–392.

Belousov, V.V., (2004), ‘Wetting of grain boundaries in ceramic materials’, Colloid Journal,66 (2): 121–127.

Bonnet, R., Cousineau, E. and Warrington, D.H., (1981), ‘Determination of near-coincidentcells for hexagonal crystals. Related DSC lattices’, Acta Crystallogr. A, 37 (2), 184–189.

Buseck, P.R., Cowley, J.M. and Eyring, L., (1988), High-Resolution Transmission ElectronMicroscopy and Associated Techniques, Oxford, Oxford University Press.

Cahn, J.W., (1977), ‘Critical point wetting’, J. Chem. Phys, 66 (8), 3667–3772.Cannon, R.M., Rühle, M., Hoffmann, M.J., French, R.H., Gu, H., Tomsia, A.P. and Saiz,

E., (2000), ‘Adsorption and wetting mechanisms at ceramic grain boundaries’, CeramicTransactions, 118, 427–444.

Cinibulk, M.K. and Kleebe, H.-J., (1993), ‘Effects of oxidation on intergranular phasesin silicon nitride ceramics’, J. Mater. Sci., 28 (21), 5775–5782.

Clarke, D.R., (1979), ‘High resolution techniques and application to nonoxide ceramics’,J. Am. Ceram. Soc., 62 (5/6), 236–246.

Clarke, D.R., (1987), ‘Grain boundaries in polycrystalline ceramics’, Ann. Rev. Mat. Sci.,17, 57–74.

Clarke, D.R., (1989), ‘High-temperature microstructure of a hot-pressed silicon nitride’,J. Am. Ceram. Soc., 72 (9), 1604–1609.

Clarke, D.R., (1990), ‘Perspectives concerning grain boundaries in ceramics’ Am. Ceram.Soc. Bull., 69 (4), 682–685.

Clarke, D.R. and Gee, M.L., (1992), ‘Wetting of surfaces and grain boundaries’, in Wolf,D. and Yip, S., (eds), Materials Interfaces, London, Chapman & Hall, 255–272.

Dietrich, S., (1991), ‘Fluid interfaces - wetting, critical adsorption, van der Waals tails,and the concept of the effective interface potential’, in Taub, H., Torzo, G., Lauter,H.J. and Fain, S.C., (eds), Phase. Transitions in Surface Films 2, NATO AdvancedScience Series, Physics, Vol. 267, 391–423.

Falk, L.K.L., (1998), ‘Electron spectroscopic imaging and fine probe EDX analysis ofliquid phase sintered ceramics’, J. Eur. Ceram. Soc., 18 (15), 2263–2279.

Flewitt, P.E.J. and Wild, R.K., (2004), Grain Boundaries: Their Microstructure andChemistry, Chichester, John Wiley & Sons.

Forwood, C.T. and Clarebrough, L.M., (1991), Electron Microscopy of Interfaces inMetals and Alloys, Bristol and Philadelphia, Institute of Physics Publishing.

French, R.H., Cannon, R.M., DeNoyer, L.K. and Chiang, Y.-M., (1995), ‘Full spectralcalculation of non-retarded Hamaker constants for ceramic systems from interbandtransition strengths’, Solid State Ionics, 75, 13–33.

Gu, H., (2002), ‘Variation of width and composition of grain-boundary film in a high-purity silicon nitride with minimal silica’, J. Am. Ceram. Soc., 85 (1), 33–37.


Gu, H., (2004), ‘Electron energy-loss spectroscopy characterization of ~1 nm-thickamorphous film at grain boundaries in Si-based ceramics’, Mater. Trans., 45 (7),2091-2098.

Gu, H., (2004), ‘Evolution of intergranular boundaries and phases in SiC and Si3N4

ceramics under high temperature deformation: case studies by analytical TEM’, Zeitschriftfür Metallkunde, 95 (4), 271-274.

Jin, Q., Wilkinson, D.S. and Weatherly, G.C., (1999), ‘High-resolution electron microscopyinvestigation of viscous flow creep in a high-purity silicon nitride’, J. Am. Ceram.Soc., 82 (6), 1492–1496.

Keyse, R.J., Goodhew, P.J., Garrattt-Reed, A.J. and Lorimer, G.W., (1998), Introductionto Scanning Transmission Electron Microscopy, Bios Scientific Publishers.

Kleebe, H.-J. and Cinibulk, M.K., (1993), ‘Transmission electron microscopycharacterization of a ceria-fluxed silicon nitride’, J. Mater. Sci. Lett., 12 (2), 70–72.

Kleebe, H.-J., Corbin, N., Willkens, C. and Rühle, M., (1990), ‘Transmission electronmicroscopy studies of silicon nitride/silicon carbide interfaces’, Mat. Res. Soc. Symp.Proc., 170, 79-84.

Kleebe, H.-J., Cinibulk, M.K., Cannon, R.M. and Rühle, M., (1993), ‘Statistical analysisof the intergranular film thickness in silicon nitride ceramics’, J. Am. Ceram. Soc., 76(8), 1969–1977.

Knowles, K.M., Turan, S., Kumar, A., Chen, S.H. and Clegg, W.C., (1999), ‘Transmissionelectron microscopy of interfaces in structural ceramic composites’, J. Microscopy,196(2) 194–202.

Krivanek, O.L., Shaw, T.M. and Thomas, G., (1979), ‘Imaging of thin intergranularphases by high-resolution electron microscopy’, J. Appl. Phys., 50 (6), 4223-4227.

Lee, W.E. and Rainforth, W.M., (1994), Ceramic Microstructures: Property Control byProcessing, London, Chapman & Hall.

Luo, J. and Chiang, Y-M., (2000), ‘Stabilization of surface films in ceramics’, CeramicTransactions, 118, 419–426.

Luo, J. and Chiang, Y-M., (2000), ‘Existence and stability of nanometer-thick disorderedfilms on oxide surfaces’, Acta Mater., 48 (18/19), 4501–4515.

More, K.L., Koester, D.A. and Davis, R.F., (1991), ‘Microstructural characterization of acreep-deformed SiC whisker-reinforced Si3N4’, Ultramicroscopy, 37, 263-278.

Ness, J.N., Stobbs, W.M. and Page, T.F., (1986), ‘A TEM Fresnel diffraction-based methodfor characterizing thin grain-boundary and interfacial films’, Phil. Mag. A, 54 (5),679-702.

Nishimura, T., Mitomo, M., Ishida, A. and Gu, H., (2000), ‘Improvement of high temperaturestrength and creep of alpha-sialon by grain boundary crystallization’, Key EngineeringMaterials, 171, 741-746.

Norton, M.G. and Carter, C.B., (1990), ‘Interfaces in structural ceramics’, MRS Bulletin,15 (10), 51-59.

Norton, M.G. and Carter, C.B., (1990), ‘Grain and interphase boundaries in ceramics andceramic composites’, in Wolf, D. and Yip, S., (eds), Materials Interfaces, London,Chapman & Hall, 151-189.

Pan, X., Gu, H., van Weeren, R., Danforth, S.C., Cannon, R.M. and Rühle, M., (1996),‘Grain-boundary microstructure and chemistry of a hot isostatically pressed high-purity silicon nitride’, J. Am. Ceram. Soc, 79 (9), 2313–2320.

Pandit, R., Schick, M., and Wortis, M., (1982), ‘Systematics of multilayer adsorptionphenomena on attractive substrates’, Phys. Rev. B, 26 (9), 5112–5140.

Pezzotti, G. and Ota, K., (1997), ‘Grain-boundary sliding in fluorine-doped silicon nitride’,J. Am. Ceram. Soc., 80 (3), 599–603.


Pezzotti, G., Tanaka, I. and Nishida, T., (1993), ‘Intrinsic fracture energy of polycrystallinesilicon nitride’, Phil. Mag. Lett., 67 (2), 95–100.

Pezzotti, G., Ota, K., Kleebe, H.-J., Okamoto, Y. and Nishida, T., (1995), ‘Viscous behaviorof interfaces in fluorine-doped Si3N4/SiC composites’, Acta. Metall. Mater., 43 (12),4357–4370.

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Reimer, L., (1997), Transmission Electron Microscopy: Physics of Image Formation andMicroanalysis, 4th edn, Berlin and Heidelberg, Springer-Verlag.

Reimer, L., Zepke, U., Moesch, J., Schulze-Hillert, S., Ross-Messemer, M., Probst, W.and Weimer, E., (1992), EELSpectroscopy, Oberkochen, Germany, Carl Zeiss.

Rühle, M., (1984), ‘TEM observations of grain boundaries in ceramics’, Mat. Res. Soc.Symp. Proc., 31, 317-323.

Sakuma, T., Sheppard, L.M., and Ikuhara, Y., (eds) (2000), Grain Boundary Engineeringin Ceramics: From Grain Boundary Phenomena to Grain Boundary Quantum Structures,Ceramic Transactions, 118, Westerville, OH, The American Ceramic Society.

Schick, M., (1990), ‘Introduction to wetting phenomena’, in Charvolin, J., Joanny, J.F.and Zinn-Justin, J., Liquids at Interfaces, Les Houches Summer School Lectures,Session XLVIII, Amsterdam, Elsevier, 415–497.

Sigl, L.S. and Kleebe, H.-J., (1993), ‘Core/rim structure of liquid-phase-sintered siliconcarbide’, J. Am. Ceram. Soc., 76 (3), 773–776.

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Smart, R., St C., and Nowotny, J., (1998), Ceramic Interfaces: Properties and Applications,London, Maney Publishing.

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18.1 Introduction

Sialons (Oyama and Kamigaito, 1971; Jack and Wilson, 1972; Hampshire etal., 1978) are a family of advanced structural ceramics that exhibit a goodcombination of properties such as high strength at elevated temperatures,dimensional stability, good corrosion and erosion behaviour, high elasticmodulus, low mass density, good thermal shock resistance and high hardness.These combined properties make them useful as structural materials. However,as with other structural ceramics, they suffer from the disadvantage of lowfracture toughness (3.0–8.0 MPa m1/2) with no R-curve behaviour (Cao andMetselaar, 1991; Shen et al., 1996a,b,c). Although they can be used forengineering components, their reliability is still not adequate in many casesfor industry to commit to production. The most attractive and effective wayof increasing toughness of sialons is through the use of a second-phasereinforcement, thus generating a sialon matrix composite. In composite form,significant improvements can be achieved such as increased fracture toughness,less strength variability, less flaw sensitivity, reduced crack propagation andbetter reliability; and even more significantly, the failure manner of sialoncomposites can be changed and controlled. Recently, several new techniqueshave been developed, such as a ¤ b sialon transformation (Mandal et a1.,1993; Thompson, 1994; Yu et al., 1998b) and the retention of elongated a-or b-sialon grains in dense a/b-sialon composites (Shen et al., 1996a; Chenand Rosenflanz, 1997; Yu et al., 2001a, b). Attempts have also been made byNordberg et al. (1993, 1997a; Nordberg and Ekström, 1995) to reinforce Y-a-sialon ceramics by incorporating a second phase, e.g. MoSi2 particles orSiC-whiskers. More recently, continuous fibre reinforcement of sialons hasbeen explored and this has proved to be an effective way of overcomingsome of the deficiencies of sialon materials. Research by Zhang and Thompson(1995, 1997), Demir and Thompson (2001) and Yu and Thompson (1998),Yu et al., 2002a, 2002b) has given very encouraging results using high-performance carbon fibres and Nicalon SiC fibres for reinforcing oxynitrideglass, a- and b-sialons.

18Sialons

Z B Y U, Queen’s University, Canada andD P T H O M P S O N, University of Newcastle, UK


Section 18.2 gives a very brief introduction to sialons. Section 18.3 outlinesthe challenges to be overcome in order to make toughened and strengthenedsialon products. Progress in developing sialon composites forms the mainpart of this chapter and is discussed in section 18.4, which deals with a/b-sialon composites, particle/whisker-reinforced sialons, and fibre-reinforcedsialons. In the final section, several conclusions and suggestions for futurework are summarised.

18.2 Sialons

Sialon is the generic name for the large family of silicon nitride solid solutionscontaining the basic elements Si, Al, O and N. Over the last three decades,the matrix sialon phases (a-, b-, and O-) have been developed and variousexcellent reviews are available (Cao and Metselaar, 1991; Ekström and Nygren,1992; Izhevskiy et al., 2000). In the sialon family, a- and b-sialons offermost interest as engineering ceramics because of their excellent combinationof mechanical and high-temperature properties.

18.2.1 b-Sialon

b-Sialon ceramics were concurrently discovered by Jack and Wilson (1972)in the UK and by Oyama and Kamigaito (1971) in Japan, who reported thatup to two-thirds of the Si and N in b-Si3N4 could be substituted by Al and Owithout change of structure to form a Si6–zAlzOzN8–z solid solution with theb-Si3N4 structure, where the z-value shows the degree of solid solubility andvaries continuously from zero to about 4.2 (Gauckler et al., 1975). b-Sialonceramics dominated the early interest because they could be pressurelesslysintered into complex shapes, and the resulting materials had a goodcombination of properties; for example, some pressurelessly sintered b-sialonceramics have strengths of up to 1000 MPa and fracture toughnesses of upto 8 MPa m1/2 (Jack, 1976).

18.2.2 a-Sialon

a-Sialon is a sialon derivative of a-Si3N4, with the general compositionMxSi12–(m+n)Alm+nOnN16–n, where, relative to the a unit cell contents of Si12N16,m (Si–N) bonds are replaced by m (Al–N) bonds and n (Si–N) bonds by n(Al–O) bonds (Ekström, 1992a,b). The charge discrepancy caused by thereplacement mechanism is compensated by the introduction of the metal ionM, with m = vx, where v is the metal ion valency. Because of the particularatomic arrangement in a-Si3N4, the unit cell, containing Si12N16, has twointerstitial sites large enough to accommodate the M atoms, and thereforex £ 2. This offers the possibility of reducing the amount of residual glassy

Sialons 493

phase if sintering additives are used which can subsequently be incorporatedinto the a-sialon structure. The existence of metal-doped a-sialons wasdemonstrated by Hampshire et al. (1978), even though the technologicalimportance of a-sialon compared with b-sialon was not appreciated at thattime. Their main advantage is that fully dense a-sialon ceramics have anexcellent hardness in excess of ~20 GPa (Nordberg et al., 1993; Shen et al.,1996a, d, e; Yu et al., 2001b).

18.2.3 a/b-Sialon and transformation

Dense b-sialon ceramics with low z values (z ~ 1) have microstructuresconsisting of elongated crystals and grain boundary glass, and of all thevarious sialon materials, show the highest observed fracture toughness valuesat room temperature (Ekström, 1989); however, their hardness is relativelylow (14–15 GPa) partly because of the residual grain boundary glass(~10 GPa). In contrast, a-sialon ceramics generally have equiaxed grainsand show an excellent hardness (~22 GPa), much higher than that of anyother sialon ceramics. The properties of mixed a/b-sialon ceramics can beoptimised by adjusting the proportion of the two phases, and this can be doneby changing the starting composition (Ekström, 1989). A high content of b-sialon yields high strength and toughness, whereas a high proportion of thea-sialon phase gives excellent hardness. Mixed a/b-sialon ceramics can beproduced with very little intergranular glass and so their wear resistance,strength and especially creep resistance at high temperatures are greatlyimproved. More recently, it has been demonstrated by Mandal et al. (1993),Thompson (1994), and Yu et al. (2000) that some densified (a + b)-sialonscan undergo in situ reversible a ¤ b sialon transformation, merely by heattreatment at temperatures below the sintering temperature. Careful control ofthe heat treatment schedule enables predetermined values of hardness, strength,and toughness to be achieved in the final (a + b) composite.

18.3 Challenges in toughening and

strengthening sialons

As with other ceramic composites, the combination of a- and/or b-sialonwith reinforcement agents results in sialon composites. This simple andobvious statement encompasses many factors which must be taken into accountfor successfully fabricating composites with a designed microstructure andimproved properties (Prewo, 1989). For sialon matrix composites, the mostimportant factors are physical compatibility including Young’s modulus, elasticstrain (Kerans and Parthasarathy, 1991) and thermal expansion coefficient(Sambell et al., 1972a, b), and chemical compatibility between sialon matrix


and reinforcement agents; the special requirements for processing andfabricating sialon composites are also a challenge.

As is well known, the development of a- and/or b-sialon composites hasnot proceeded quickly, even though some of the earliest ceramic matrixcomposites investigated were based on a Si3N4 matrix (Guo et al., 1982).The difficulties in achieving a fully dense sialon composite at relative lowtemperatures are greater than previously encountered for either glass or glass–ceramic materials because (1) the formation temperature of sialons isapproximately 1500∞C (Kuang et al., 1990; Yu and Thompson, 1998); (2) thedensification temperature of sialons is higher than 1650∞C (Cao and Metselaar,1991); and (3) liquid phases formed during the formation and densificationof sialons are substantially accommodated into the sialon structure as sinteringproceeds, especially for a-sialons, thereby making final densification difficult.Therefore, formation and densification of sialons are not easy, and highdensities cannot be easily achieved at low temperatures, short sintering times,and under pressureless sintering conditions; Furthermore, the incorporationof reinforcement agents, especially fibres, into a sialon matrix makes thedensification even more difficult, and the high hot-pressing temperatures andpressures needed inevitably tend to create a very strong bond between thesialon matrix and the reinforcement agents or to degrade the reinforcementagents. Therefore, all these factors make it exceptionally difficult to achievea fully dense sialon composite.

18.4 Sialon composites

Over the last decade, considerable efforts have been committed to thetoughening of sialons and substantial progress has been achieved using variousreinforcements. According to the form of reinforcement, sialon compositescan be classified as either particle reinforced, discontinuous fibre (whiskers/short fibres) reinforced, or continuous fibre reinforced.

18.4.1 a /b-Sialon composites

Toughening mechanisms in a § b-sialon composites are similar to thoseoperative in second-phase particle reinforced composites, but, rather thanthe deliberate addition of a second phase, a§b-sialon composites are fabricatedby simultaneous crystallisation of the two solid solutions a- and b-sialonfrom a eutectic composition liquid. This requires careful design of the startingcomposition which is usually located within the (a + b)-sialon region of thea-sialon plane as illustrated in Fig. 18.1.

The a/b phase ratio can be controlled just by changing the overallcomposition in most sialon systems. b-Sialon grains normally grow in elongatedshape with a high aspect ratio (Lange, 1979; Wötting et al., 1986), and this

Sialons 495

promotes mechanisms such as crack bridging and deflection. a-Sialon withvery high hardness used to be considered as always occurring with equiaxedgrains and thus its fracture toughness is inferior to that of b-sialon. Theseobservations led to work programmes being focused on the design of a§b-sialon composites that combine the strength and fracture toughness of b-sialon and the hardness of a-sialon to give improved mechanical properties.For example, it has been reported that the composition containing 50% b-sialon and 50% a-sialon showed a hardness (HV10) of ~22 GPa and a fracturetoughness of 5.5 MPa m1/2 (Ekström, 1997). Jones et al. (2003) chose Y-a§b-sialon composites with compositions lying on the Si3N4–9AlN.Y2O3 lineand produced sialon composites with different a/b-sialon ratios (75–51%),the bimodal microstructures of the resulting sialon composites showingimproved wear properties compared to their monolithic counterparts. Morerecently, it has been observed that dense a-sialons containing elongatedgrains can also be prepared by carefully selecting the starting composition(Shen et al., 1996a; Chen and Rosenflanz, 1997; Nordberg et al., 1997b,Wood et al., 1999; Yu et al., 2001a); these are encouraging results and showthe good potential for obtaining materials with high hardness and toughness.Therefore, a/b composites have become an important research area, anddetails of the starting powder, composition, processing route, etc., should becarefully considered in order to obtain the desired phases, grain shape andsize to optimise the resulting bimodal microstructure.

Composition design

Different starting powders and chemical compositions result in differentphases, grain boundary states, and microstructures (see Table 18.1). For

a-sialon region

a + b-sialon region

43

(Al2O3·AIN)

(M2O3·9AIN)

MN.3AIN

b-sialon line n = 3.0

n = 2.0

n = 1.0

Si3N4

m = 1.0

m = 2.0

m = 3.0

18.1 Schematic illustration of a- and b-sialon phase regions on the a-sialon plane (M: metal cation with a valence of 3+).


instance on the a-sialon plane, for a general a-sialon compositionMxSi12–(m+n)Alm+nOnN16–n, designated as (m, n), the microstructures of boththe low (m, n) composition (1.0, 1.0) and the high (m, n) composition (2.0,2.0) consist mainly of equiaxed a-sialon grains, whereas (0.5, 2.0) and (1.0,2.0) results show a bimodal microstructure consisting of equiaxed and elongatedgrains of the two main sialon phases. Also, samples prepared with high b-phase content silicon nitride powder show a lower toughness than thoseprepared from a high a-phase starting silicon nitride powder.

Volume fraction of elongated grains

According to the results of Faber and Evans (1983a,b), the majority of thetoughening from crack deflection develops for a volume fraction ofreinforcement of <20% for a given morphology of the reinforcing phase.Figure 18.2 shows typical microstructures for sialon samples with differentcompositions. The volume of the elongated grains in the microstructureshown in Fig. 18.2(b) is over 20%, and an excess of elongated grain incomposition (0.5, 2.0) does not add to the fracture toughness. So, it might beexpected that the contribution of the volume fraction of elongated grains inthese sialons to crack deflection would be comparable. However, as Table18.1 shows, the (0.5, 2.0) composition gives an approximately 50% higherfracture toughness compared to the (1.0, 2.0) sample; this can be attributedto other factors as discussed below.

Table 18.1 Effects of starting powder, composition on phase present,microstructure and mechanical properties of a-sialons

Composition Starting a § b Sialon Morphology HV10 KIc

(m, n) powder* ratio of grains (GPa) (MPa m1/2)

(0.5, 2.0) a 41/59 Equiaxed 17.9 ± 0.5 6.8 ± 0.5+ elongated

(1.0, 1.0) a 100/0 Equiaxed 20.3 ± 0.4 4.9 ± 0.3

(1.0, 1.0) b 100/0 Equiaxed 20.2 ± 0.9 4.2 ± 0.4

(1.0, 2.0) a 100/0 Equiaxed 18.3 ± 0.5 5.0 ± 0.4+ elongated

(1.0, 2.0) b 100/0 Equiaxed 18.0 ± 0.2 4.4 ± 0.3+ elongated

(2.0, 2.0) a 100/0 Equiaxed 15.3 ± 0.7 4.7 ± 0.7

*a and b stand for a- and b-Si3N4 respectively.Reprinted from J. Mater. Sci., 36(14), Yu Z B, Thompson D P and Bhatti A R, ‘Self-reinforcement in Li-a-sialon ceramics’, Fig. 2, 3343–3353, (2001). Copyright 2001,with kind permission of Kluwer Academic Publishers.

Sialons 497

The aspect ratio effect

The length:diameter (L/D) ratio of the elongated grain plays an importantrole in toughening. The average L/D of elongated sialon grains in sample(0.5, 2.0) is about 10–12, which is higher than that in the (1.0, 2.0) samplewith an average L/D of ~6. It was predicted (Shalek et al., 1986) that rod-shaped obstacles with large aspect ratios should impart maximum toughnessbecause high aspect ratios result in a large twist angle of the crack. Themaximum effect which can be achieved with rod-like particles is:

Gc/Gm ª 4.0 (when L/D = 12) (18.1)

Gc/Gm ª 3.0 (when L/D = 3) (18.2)

where Gc = crack resistance force of the composite, and Gm = crack resistanceforce of the matrix. Therefore, the contribution of the elongated sialon grainsin (0.5, 2.0) samples towards the fracture toughness is ~1.2 times higher thanthat of elongated sialon grains in (1.0, 2.0) samples as deduced from theL/D ratio.

Thermal and Young’s modulus mismatches

Elongated a- and b-sialon grains have different thermal expansion coefficientsand slightly different Young’s moduli, which cause local stress fields to formaround the elongated grains, and these play a large role in the crack deflectionprocess. Compared to pure (1.0, 2.0) a-sialon, the multiphase a § b-sialon ofcomposition (0.5, 2.0) showed higher fracture toughness. This can be partiallyexplained by the physical mismatches. Assuming plane strain and isotropicelastic behaviour, the thermal residual stresses P are proportional to the

(a) (b)

18.2 (a) Microstructures of sialons with different compositions: (a)(1.0, 1.0), (b) (0.5, 2.0) (reprinted from J. Eur. Ceram. Soc., 21(13), YuZ B, Thompson D P and Bhatti A R, ‘In-situ growth of elongatedgrains in Li-a-sialon ceramics’, 2423–2434 (2001). Copyright 2001,with permission of Elsevier).


difference between the thermal expansion coefficients of the equiaxed andelongated sialon grains estimated according to the analysis of Budiansky etal. (1986) and Yu et al. (2001b):

P = E feq

1

(1 – )

2 (1 – )

el n a

(18.3)

wheree = (ael – aeq )DTl1 = 1 – (1 – 2n)[1 – f + (1 – f )Eeq/Eel]/[2(1 – n)]f = volume fraction of the elongated grainsael, aeq = linear thermal expansion coefficientsEeq, Eel = Young’s modulineq, nel = Poisson’s ratiossubscripts el and eq refer to the elongated and equiaxed sialon grains,respectively, and here it is assumed that neq = nel = n.

The residual stress plays an important role in the deflection process sinceit actually determines the type of crack–microstructure interactions. The(1.0, 2.0) sample consists of pure a-sialon reinforced with in situ formedelongated a-sialon grains, and the local residual stress induced by thermaland elastic mismatch around the elongated a-sialon grains is nearly zero, i.e.aela – aeqa = 0 and Eela – Eeqa = 0. Therefore, the residual stress makes littlecontribution to the crack deflection, and the crack tilting and twisting isrelatively small because they are only caused by the elongated grains. However,in the (0.5, 2.0) composite consisting of a-sialon grains reinforced with insitu formed elongated b-sialon needles, there exist local residual stressesaround the elongated b-sialon grains induced by thermal and elastic mismatch,because the thermal expansion coefficients and Young’s moduli between thematrix, the equiaxed a-sialon grains (aeq, Eeq) and the elongated grains (aelb,Eelb) are slightly different, i.e. aelb – aeqa π 0, and Eelb – Eeqa π 0. Therefore,the residual stress field in sample (0.5, 2.0) makes an additional contributionto crack deflection. So the crack tilting and twisting are more marked in the(0.5, 2.0) sample than in the (1.0, 2.0) sample and the fracture toughnessreflects this.

18.4.2 Particle-reinforced sialons

Because conventional processing routes can be used, particle reinforcementis a more attractive fabrication route for making a§b-sialon composites, andoffers other advantages, such as cheapness and the avoidance of health hazardsassociated with the use of whiskers or short fibres. Apart from a§b-sialoncomposites, various kinds of particles including ZrO2, SiC, TiN, and MoSi2,among others, have been suggested and explored as the reinforcing phase.

Sialons 499

Toughening mechanisms using ZrO2 particles in silicon nitride or sialon-based matrices can be tailored to take advantage of the phase transformationand this is still a popular research area (Ekström et al., 1991; Yu et al.,2002a). A slight increase in both strength and toughness can be observed forboth silicon nitride and sialons. However, ZrO2 and sialons are not chemicallycompatible and there is some difficulty in controlling the reactions betweenthem at sintering temperatures.

SiC particles have been used as reinforcements for Si3N4 and sialon ceramicsbecause of the good compatibility between nitrides and SiC. Buljan et al.(1987) found that an increase in toughness was observed for coarse ratherthan fine SiC particles since the coarser particles can promote large crackdeflection. Niihara et al. (1990) reported that Si3N4/32vol%SiC nanocompositeswith 8 wt% of Y2O3 sintering additive exhibited strengths of 970 MPa and840 MPa at 1400∞C and 1500∞C, respectively, and indicated that the nano-size SiC particles were mainly dispersed within Si3N4 grains with littleadditional grain boundary material. More recently, the preparation of b-sialon composites reinforced by nano-SiC particles generated by in situchemical techniques has been explored by Li et al. (1997), who found thatthe chemical route is favourable for the sintering of the composite. Liu et al.(1999) demonstrated that when 10–20 wt% SiC nanoparticles were added toa series of rare-earth doped Ln-a-sialons (Ln = Nd, Dy, Yb) of compositionLn0.33Si9.3Al2.7O1.7N14.3, hot-pressed at 1800∞C and then heat-treated at 1450∞Cto modify the microstructure and properties by a § b-sialon transformation,an enhanced hardness was found in proportion to the content of SiC particlesand with an increase in the a:b-sialon ratio, with a maximum hardness of 22GPa (HV10) being achieved; however, a decrease in fracture toughness wasobserved. It was suggested that the significant increase in high temperaturestrength could be attributed to residual microstresses around the SiC particlesdue to the difference in thermal expansion coefficients between Si3N4/sialonsand SiC. The strengthening and toughening mechanism of nano-SiC particlereinforced sialon ceramics has generated many current research programmes.MoSi2 particulate reinforced a-sialon composites have been studied byNordberg and Ekström (1995) who found that the hardness of fully densesamples changed from 22.5 to 15.3 GPa with KIc increasing from 3.4 to 5.2MPa m1/2, when up to 30 vol% MoSi2 was present. Particle-reinforced sialoncomposites are a new direction and some progress has been made; however,to achieve a significant toughening effect, more work is needed in this area.

18.4.3 Whisker-reinforced sialons

Whiskers provide a better reinforcement than particles from the point ofview of the effect of dispersoid shape on toughness based on theoreticalpredictions (Shalek et al., 1986). The most commonly used whisker


reinforcement for sialons is SiC. SiC whiskers have been widely used intoughening and strengthening Si3N4 ceramics (Buljan et al., 1987; Campbellet al., 1990) and it has been observed that the strength and fracture toughnessincreased by 25% and 40% respectively, compared with monolithic Si3N4

again depending on sintering additives, sintering temperatures and whiskertype. The incorporation of SiC whiskers into Y-doped a-sialon ceramics hasbeen explored by Nordberg et al. (1993), for composites containing 30 wt%SiC whiskers. Unfortunately, only a minor part of the SiC whisker/sialonmatrix interfaces consisted of a 4–5 nm thick amorphous layer (which facilitatedpullout), whereas a direct and strong bond was the most commonly occurringcontact. The hardness of the composites increased to ~21 GPa, but the fracturetoughness increased only slightly because of the strong bonding betweenwhiskers and matrix.

Although fracture toughness can be increased, particle- or whisker-reinforcedsialon composites generally show brittle behaviour and low damage tolerance;this is in contrast to fibre-reinforced sialons which exhibit non-catastrophicfailure.

18.4.4 Carbon fibre (Cf)-reinforced sialons

Carbon fibres offer an excellent combination of strength, modulus, andtoughness compared with all other reinforcing fibres, and among availablematerials, carbon fibres can retain good mechanical properties above 1000∞C.Carbon fibres have been widely used in ceramic composite systems (Guo etal., 1982; Prewo, 1982; Phillips et al., 1990; Zhang and Thompson, 1995,1997; Yu and Thompson, 1998; Yu et al., 2002a). Yu and Thompson (1998)demonstrated a new two-step process in developing carbon fibre-reinforceda-sialons to deal with the problems encountered in fabricating sialoncomposites. In their process, first of all, a-sialon was pre-prepared by pressure-less sintering and a lithium aluminium oxide glass (LAS) sintering additivewas prepared by conventional melting followed by crushing into a fine powder.Second, continuous fibre-reinforced sialon matrix composites, with the as-prepared a-sialon powder and sintering additives (present as 25% by volumeof the matrix phase), were produced using slurry infiltration followed by hot-pressing and achieved apparent densities of up to 96% of theoretical densitybelow 1550∞C for composites containing as much as 60 vol% of carbonfibres. In the next section, effects such as thermal mismatch, fibre distributionand processing parameters in carbon fibre-reinforced sialons are discussed.

Effects of processing parameters on Cf /sialon composites

As shown in Table 18.2, both the volume content of Cf and the sinteringtemperature affect the properties and fracture behaviour of Cf /a-sialon

Sialons 501

composites. The typical load–displacement curves and fracture surfaces forthese composites are given in Fig. 18.3.

Below 1500∞C, a typical load–displacement curve for a Cf/sialon compositeindicates a characteristic first matrix cracking event followed by a region ofintermittent cracking non-linear behaviour until eventual failure occurs atsome maximum load. The load drop occurs at a characteristic value followedby a long ‘tail’. Fracture at this maximum load is generally associated withfibre bundle fracture, and the tail region of intermittent cracking is associatedwith fibre bridging and pullout (Evans, 1985; Evans et al., 1991). Cracking,debonding and pullout are the main toughening processes operating in thesematerials during failure. However, sialon composites sintered at hightemperatures (e.g. 1550∞C) give poor mechanical performance (Table 18.2)

Table 18.2 Effect of sintering temperature on the properties of Cf/sialon composites

Sample Properties 1350∞C 1400∞C 1450∞C 1500∞C 1550∞C

a + 60 KIc — 7.3 ± 1.0 12.7 ± 2.0 9.4 ± 2.0 3.6 ± 1.0vol% (Cf) (MPa m1/2)

a + 60 s (MPa) — 165 ± 10 223 ± 20 224 ± 20 93 ± 10vol% (Cf)

a + 40 s (MPa) 189 ± 10 — 337 ± 20 — 147 ± 10vol% (Cf)

Reprinted from J. Eur. Ceram. Soc., 22(2), Yu Z B, Thompson D P and Bhatti A R,‘Synergistic roles of carbon fibres and ZrO2 particles in strengthening andtoughening Li-a-sialon composites’, 225–235 (2002). Copyright 2002, withpermission of Elsevier.

(b)

Displacement (mm)

(a)

200

150

100

50

0

Load

(N

)

0.02

0.10

0.17

0.24

0.31

0.42

0.78

1.14

1.50

1.86

18.3 Fracture behaviour of Cf/a-sialon hot-pressed at 1500∞C: (a)load–displacement curve; (b) fracture surface (reprinted from J. Eur.Ceram. Soc., 22(2), Yu Z B, Thompson D P and Bhatti A R,‘Synergistic roles of carbon fibres and ZrO2 particles in strengtheningand toughening Li-a-sialon composites’, 225–235 (2002). Copyright2002, with permission of Elsevier).


and usually show a linear load–displacement behaviour up to a maximumvalue followed by a sudden decrease in load with a very limited continueddeformation, suggesting that serious degradation of the carbon fibre occursat this sintering temperature. A significant improvement in fracture toughnesscan be achieved in Cf/a-sialon ceramics sintered at 1450∞C, with KIc valuesas high as 12.7 MPa m1/2. However, the bending strength is low (~222 MPa),which can be in part attributed to thermal mismatch between the matrix andthe fibre and also to non-homogeneous distribution of the fibres, which isoften found in carbon fibre-reinforced ceramic matrix composites.

Thermal mismatch in Cf /a-sialon composites

The properties of fibrous composite materials are strongly dependent onmicrostructural parameters such as fibre diameter, fibre length, fibre distribution,volume fraction of the fibres and the alignment and packing arrangement ofthe fibres. It is important to control these parameters for effective design andmanufacture of these composites. An ideal situation is that all the fibresshould be aligned parallel to each other and distributed homogeneously, butthis is not easy to do in practice. An example is given in Fig. 18.4, whichshows micro-cracks occurring in Cf/a-sialon composites containing ~60 vol%of carbon fibres induced by non-homogeneous distribution of the fibres andthe thermal mismatch between Cf and the sialon matrix.

18.4 Micro-crack induced by non-uniform distribution of carbonfibres and thermal mismatch between Cf and sialon in the Cf /a-sialoncomposite (reprinted from J. Eur. Ceram. Soc., 22(2), Yu Z B,Thompson D P and Bhatti A R, ‘Synergistic roles of carbon fibres andZrO2 particles in strengthening and toughening Li-a-sialoncomposites’, 225–235 (2002). Copyright 2002, with permission ofElsevier).

Sialons 503

Assuming that continuous carbon fibre reinforced a-sialon compositesare composed of two materials, namely carbon fibre layers and a/LAS matrixlayers, forming alternate sheets with linear thermal expansion coefficientsaf, am, elastic moduli Ef, Em, and Poisson’s ratios nf, nm, then the boundarystresses in the laminate can be estimated in the same way as given by Kingeryet al. (1976). When the composite is sintered at a relatively high temperature(Ts ~ 1450∞C), and in a stress-free condition, and is cooled down to roomtemperature Tr, the temperature difference DT = Tc – Tr,, where Tc is thetemperature below which the matrix ceases to behave viscoplastically, willresult in thermal expansion changes in the reinforcing carbon fibres of afDTand in the a-sialon/LAS matrix of amDT. These expansions are not the sameand the composite must adopt an intermediate overall expansion, determinedby the relative elastic moduli and volume fractions of fibres and matrix, butwith the net compressive force on the carbon fibres equal to the net tensileforce in the a-sialon matrix. If s is this stress, V the volume fraction, and ethe actual strain, then when the composite is cooled down to room temperature,for a composite in which the continuous fibres are distributed homogeneously(Kingery et al., 1976):

smVm + sfVf = 0 (18.4)

Based on this model, the thermal stress in the matrix has been deduced by Yuet al. (2002a) to be:

s n Da Dmm f

ff =

– 1 –

E EV TÊ

Ëˆ¯ (18.5)

From equation (18.5), the stress caused by thermal mismatch in the matrix isproportional to the volume fraction of the fibres. Therefore, large thermalstresses in some local regions, where the distribution of carbon fibres is notuniform, will result in some cracks in the composite.

It is worth pointing out that carbon fibre itself has anisotropic thermalexpansion properties, and therefore this mismatch between the carbon fibresand the a-sialon matrix should be considered in both the radial and axialdirections when carbon fibres are unidirectionally aligned in the composite.The thermal stress caused by thermal expansion differences between thecarbon fibres and the matrix in the radial (sr) and axial (sa) directions can beestimated from the formulae (Chawla, 1993; Kerans and Parthasarathy, 1991):

sa a D

n nrf m f r m

f m m f =

– [ – ) + / ](1 + ) + (1 – )

qE E r T A rE E

◊(18.6)

sa = (am – af·a)DTE V

V E Ef f

f f m[ / ) – 1] + 1(18.7)


whereDT = temperature change during coolingq = an adjustable parameter (normally taken as unity)af.r, af.a = thermal expansion coefficients of the fibres in radial and axialdirections respectivelyam = matrix thermal expansion coefficientA = amplitude of fibre roughnessr = fibre radiusEf, Em = elastic modulinf, nm = Poisson’s ratiossubscripts f and m refer to the fibre and matrix, respectively.In Cf/sialon composites, because af.r (~8 ¥ 10–6 ∞C–1) is larger than am

(~5.6 ¥ 10–6 ∞C–1) in the radial direction, the thermal mismatch betweenfibres and matrix is acceptable. However, in the axial direction, af.a (~0) ismuch smaller than am, and on cooling from the hot-pressing temperaturewill put the matrix in tension, generating stresses of ~466 MPa and 700 MPafor a Cf/sialon composite containing 40 vol% and 60 vol% of carbon fibresrespectively (Yu et al., 2002a), which in both cases is greater than the strengthof the matrix and therefore results in the appearance of transverse cracks.

Compensation in thermal mismatch

To reduce residual stresses, careful control of the slurry infiltration processingparameters is necessary, but this may not be enough. Alternative proceduressuch as incorporating additions of ZrO2 (Guo et al., 1982; Zhang and Thompson,1995; Yu et al., 2002a) as well as introducing other phases into the matrixhave been explored.

By making use of the 3–5% volume increase accompanying the t Æ mZrO2 transformation, Yu et al. (2002a) demonstrated that addition of ZrO2 toa sialon matrix is an effective route for compensating for the thermal expansiondifferences in the matrix. As shown in Fig. 18.5, for Cf/sialon composites,when 20 wt% ZrO2 was added to the sialon matrix, all microcracks wereeliminated. When these cracks were eliminated, the strength and fracturetoughness of the Cf/a-sialon composites were significantly improved (seeTable 18.3). The KIc value for carbon fibre-reinforced a-sialons containing20 wt% ZrO2 reached a value of ~12 MPa m1/2, with a bending strength of~410 MPa, which is an excellent combination of these key properties.

The introduction of ZrO2 plays an important role in changing the thermalbehaviour of the composite. However, whether the t Æ m ZrO2 phasetransformation occurs or not is conditional on other factors, such as theparticle size of ZrO2, chemical composition and the local stress situation.Briefly, it is important to ensure that the t Æ m transformation successfullyoccurs immediately below the softening temperature of the glass during

Sialons 505

cooling, and that the amount of ZrO2 addition required is not unreasonable(Yu et al., 2002a).

18.4.5 SiC fibre-reinforced sialons

SiC fibres (SiCf) are some of the more refractory ceramic fibres, and performvery satisfactorily in oxidising environments. A key problem in the developmentof SiC fibre composites is thermal degradation of the SiC fibre (Johnson etal., 1987). The properties of SiC fibre start degrading above 600∞C because

18.5 Back-scattered SEM photo showing the elimination ofmicrocracks in Cf/a-sialon composites with the addition of 20 wt%ZrO2 (black strips: carbon fibres; white dots: ZrO2) (reprinted from J.Eur. Ceram. Soc., 22(2), Yu, Z B, Thompson D P and Bhatti A R,‘Synergistic roles of carbon fibres and ZrO2 particles in strengtheningand toughening Li-a-sialon composites’, 225–235 (2002). Copyright2002, with permission of Elsevier).

Table 18.3 Properties of carbon fibre/a-sialon composites with and without addedZrO2

Temperature ZrO2 content Strength Toughness Young’s modulus(∞C) (wt%) (MPa) (MPa m1/2) (GPa)

1400 0 190 ± 10 7.0 ± 1.0 —20 336 ± 10 12.5 ± 1.5 171.4

1450 0 218 ± 5 5.6 ± 0.5 —20 410 ± 10 11.3 ± 1.0 182.0

1500 0 105 ± 5 4.0 ± 0.2 —20 120 ± 10 3.5 ± 0.2 180.0

(Reprinted from J. Eur. Ceram. Soc, 22(2), Yu Z B, Thompson D P and Bhatti A R,‘Synergistic roles of carbon fibres and ZrO2 particles in strengthening andtoughening Li-a-sialon composites’, 225–235 (2002). Copyright 2002, withpermission of Elsevier).


of the thermodynamic instability of the composition and the microstructure(Chawla, 1993). Another problem is chemical reactions with the matrix. Inthis section, new processing techniques such as heat-treatment of SiC fibresand low-temperature processing routes involving selection and introductionof low-temperature sintering additives combined with careful control of processparameters, are shown to be useful for successful fabrication of a- and b-sialon composites.

Effect of heat-treatment of SiC fibres on SiCf /sialon composites

The effect of high-temperature heat-treatment on SiC fibre strength has beenexamined for a variety of environments (Mah et al., 1984; Pysher et al, 1989;Bibbo et al., 1991; Bodet and Tressler, 1995). More recently, Demir andThompson (2001) showed that SiC fibres heat-treated at 1600∞C for 30 minunder 45 bar of CO atmosphere gained 50% in strength while at the sametime a thin carbon coating was deposited on the fibre. These changes offersignificant advantages for the use of SiC fibres to reinforce sialons at hightemperatures. These authors used the as-received Nicalon SiC fibres (NL-207) and then the heat-treated Nicalon SiC fibres as reinforcements for thestrengthening and toughening of b-sialons at the z = 1 and z = 2 compositions,

Table 18.4 Mechanical properties of as-received and heat-treated SiCf /b-sialoncomposites

b-Sialon Fibre* Sintering Temperature TD Strength Toughnessz-value additives (∞C) (%) (MPa) (MPa m1/2)

1 A MgO, Sm2O3 1600 91 349 ± 31 11.71 ± 1.8

1 H MgO, Sm2O3 1600 98 687 ± 52 11.04 ± 1.5

1 A MgO, Y2O3, 1550 90 375 ± 112 10.66 ± 2.1Sm2O3

1 H MgO, Y2O3, 1550 95 623 ± 53 13.36 ± 1.8Sm2O3

3 A MgO, Y2O3, 1550 98 419 ± 55 11.93 ± 1.9Sm2O3

3 H MgO, Y2O3, 1550 99 627 ± 25 11.87 ± 0.7Sm2O3

3 A Li2O, MgO, 1450 94 408 ± 32 11.03 ± 2.5SiO2

3 H Li2O, MgO, 1450 98 512 ± 24 10.50 ± 0.4SiO2

*A: as-received fibre; H: heat-treated fibreReprinted from J. Eur. Ceram. Soc., 21(5), Demir A and Thompson D P, ‘High-performance SiC-fibre reinforced b-sialon CMCs prepared from heat-treatedNicalon fibres’, 639–647 (2001). Copyright 2001, with permission of Elsevier.

Sialons 507

and found a marked difference in mechanical properties of hot-pressed samplesfor the as-received and heat-treated fibres as summarised in Table 18.4.Table 18.4 shows that in the selected compositions all the samples show highfracture toughness, whereas for bending strength most samples show anenormous strength increase for the heat-treated fibres, demonstrating thatthe latter achieved toughening and strengthening simultaneously. In addition,it seems that the heat-treated fibres also played a role in facilitating thedensification of the sialon composites; moreover, heat-treatment of the fibresallows high-temperature processing of samples.

Effect of sintering additives on SiCf /sialon composites

Sintering additives have a strong effect on the interface between matrix andfibres. As shown in Table 18.5, composites hot-pressed with different sinteringadditives – lithium aluminosilicate glass (LAS), nitrogen-containing lithiumaluminosilicate glass (NLAS), and mixed Y2O3 + Al2O3 + ZrO2 (YAZ) additives– show different fibre pullout effects. Fibre pullout can be achieved incomposites with LAS and NLAS as sintering additives sintered at lowtemperatures and but not with the YAZ additive. The microstructures offracture surfaces of the SiCf/a-sialon composites with different sinteringadditives are shown in Fig. 18.6. The compositional change from the matrixto the fibre centre in SiCf/sialon composites with different sintering additivesis shown in the form of the X-ray line scan (Fig. 18.7).

YAZ as a sintering additive causes some chemical reaction at the interface

Table 18.5 Effect of sintering conditions and additives on fibre pullout inSiCf /a-sialons

Fibre Matrix/sintering Hot-pressing Maximumadditives conditions fibre pullout

T (∞C) P (MPa) t (min)length (mm)

SiC a § LAS 1400 15 10 ~251450 15 10 ~51500 15 10 ~0

SiC a /NLAS 1350 15 10 ~801400 15 10 ~401450 15 10 ~10

SiC a /YAZ 1400 15 10 ~01450 15 10 ~01500 15 10 ~0

Reprinted from Composites: Part A, 33(5),Yu Z B, Thompson D P andBhatti A R, ‘Fabrication and characterisation of SiC fibre reinforcedlithium-a-sialon matrix composites’, 621–629 (2002). Copyright 2002, withpermission of Elsevier.


between the fibre and the matrix. Aluminium and oxygen from the matrixhave diffused into the fibre, and silicon has diffused out of the fibre, forminga reaction layer about 1.5 mm in thickness, which results in such a stronginterface that neither debonding nor pullout can take place during crackpropagation in the composite. When LAS was used as the sintering additive(Fig. 18.7(b)), there was no obvious chemical reaction between the fibre andthe matrix, which is further confirmed by the X-ray line scan result. WhenNLAS was used as the sintering additive, the interface between the fibresand the matrix was very clear (Fig. 18.6(c)) and there is no obvious chemicalreaction on the surface (Fig. 18.7(c)), showing good chemical compatibilitybetween fibre and matrix.

Residual stresses in SiCf /a-sialon composites

Thermal mismatch between SiC fibres and a sialon matrix is also an issuewhich can cause poor mechanical performance, even though there may beexcellent chemical and physical compatibility between the fibres and the

18.6 Micrographs of fracture surfaces of continuous SiCf/a-sialoncomposites sintered with (a) YAZ, (b) LAS, and (c) NLAS additives(reprinted from Composites: Part A, 33(5), Yu Z B, Thompson D P andBhatti A R, ‘Fabrication and characterisation of SiC fibre reinforcedlithium-a-sialon matrix composites’, 621–629 (2002). Copyright 2002,with permission of Elsevier).

(b)(a)

(c)

Sialons 509

matrix. For instance, in SiCf/a-sialon composites with LAS as the sinteringadditive, the coefficient of thermal expansion of the SiC fibre, aSiCf (3.4 ¥10–6∞C–1), is lower than that of the a /LAS matrix (~5.6 ¥ 10–6∞C–1). Duringcooling from high temperatures, radial compressive stresses build up becausethe matrix shrinks more round the fibres. The magnitude of this stress can beapproximately calculated from equation (18.6).

A calculated radial compressive stress of 246 MPa on the fibre(corresponding to a frictional stress as large as 24.6 MPa) was obtained byYu et al. (2002b), which was big enough to cause the SiC fibres to breakwhen they were torn from the a-sialon matrix.

Therefore, in SiC fibre-reinforced sialon composites, thermal treatmentof fibres, thermal mismatch and chemical reactions between the fibre and thesialon matrix are significant factors affecting the interfacial bonding in thesecomposites. The sintering additive plays an important role in controlling thenature of the interface and requires careful selection.

18.7 X-ray line scan of O, Al and Si in continuous SiCf/a-sialoncomposites sintered with (a) YAZ, (b) LAS, and (c) NLAS additivesrespectively (reprinted from Composites: Part A, 33(5), Yu Z B,Thompson D P and Bhatti A R, ‘Fabrication and characterisation ofSiC fibre reinforced lithium-a-sialon matrix composites’, 621–629(2002). Copyright 2002, with permission of Elsevier).

(a) (b)

(c)

SiKa, 276

Alka, 63

OKa, 13

OKa, 13

Alka, 91

SiKa, 563

OKa, 12

Alka, 41

SiKa, 201


18.5 Future trends

Sialon composites can be produced by in situ growth of elongated a- or b-sialon grains, providing a self-reinforcement mechanism involving crackdeflection and bridging. The reinforcement effect is mainly determined bythe overall composition which in turn determines the morphology and thetype of matrix sialon phase present; a § b-sialon matrix offers the best optionfor producing a final composite with the desired microstructures and properties.Sialons reinforced by a secondary particulate phase or short fibres/whiskersgives some toughening effect. The most marked strengthening and tougheningeffects (achieved simultaneously) in sialons are obtained using fibre-reinforcedcomposites, and these show controlled fractured behaviour and high damagetolerance. Avoiding physical and chemical incompatibilities between thesialon matrix and the fibres is very important, and initial calculations areessential to determine whether it is necessary to tailor the properties of thesialon matrix or to modify the properties of the fibres or both.

There is also scope for the development of new techniques such as chemicalvapour infiltration (CVI) (Caputo and Lackey, 1984; Caputo et al., 1985),normal chemical reaction bonding processes, laminar sialon composites, etc.More recently, laminated composites in non-oxide and sialons havedemonstrated very promising results for strengthening (Goto and Kato, 1998)and even achieved a non-brittle failure behaviour accompanied by high damagetolerance (Yu and Krstic, 2003; Yu et al., 2005).

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Johnson, L.F., Hasselman, D.P.H. and Minford, E., (1987), ‘Thermal conductivity ofcarbon fibre-reinforced borosilicate glass’, J. Mater. Sci., 22, 3111–3117.

Jones, M.I., Hirao, K., Hyuga, H., Yamauchi, Y. and Kanzaki, S., (2003), ‘Wear propertiesof Y-a /b composite sialon ceramics’, J. Eur. Ceram. Soc., 23(10), 1730–1750.

Kerans, R.J. and Parthasarathy, T.A., (1991), ‘Theoretical analysis of the fibre pullout andpush out test’, J. Am. Ceram. Soc., 74(7), 1585–1596.

Kingery, W.D., Bowen, H.K. and Uhlmann, D.R., (1976), Introduction to Ceramics, NewYork, John Wiley & Sons.


Kuang, S.F., Huang, Z.K., Sun, W.Y., and Yen, T.S., (1990), ‘Phase relationships in theLi2O-Si3N4-AIN system’, J. Mater. Sci Lett. (9), 72–74.

Lange, F.F., (1979), ‘Fracture toughness of Si3N4 as a function of the initial alpha phasecontent’, J. Am. Ceram. Soc. 62(7–8), 428–430.

Li, Q., Gao, L., Jiang, D., Zhang, C. and Yan, D.S., (1997), ‘Preparation of b-SiAlON/nano-SiC composites’, J. Mater. Sci. Lett., 16, 1620–1621.

Liu, Q., Gao, L., Yan, D.S. and Thompson, D.P., (1999), ‘Hard sialon ceramics reinforcedwith SiC nanoparticles’, Mater. Sci. Eng., A269, 1–7.

Mah, T., Hecht, N.L., McCullen, D.E., Honigman, J.R., Kim, H.M., Katz, A.P. andLipsitt, H.A., (1984), ‘Thermal stability of SiC fibre (Nicalon)’, J. Mater. Sci., 19(4),1191–1201.

Mandal H., Thompson D.P. and Ekström T. (1993), ‘a ¤ b Transformation in heat-treated sialon ceramics’, J. Eur. Ceram. Soc., 12(6), 421–429.

Niihara, K., Izaki, K. and Kawakami, T., (1990), ‘Hot-pressed Si3N4–32%SiC nanocompositefrom amorphous Si–C–N powder with improved strength above 1200∞C’, J. Mater.Sci. Lett., 10, 112–114.

Nordberg, L.O. and Ekström, T., (1995), ‘Hot-pressed MoSi2-particulate-reinforced a-SiAlON composites’, J. Am. Ceram. Soc., 78(3), 797–800.

Nordberg, L.D., Esktröm, T. and Wen, S., (1993), ‘Hot-pressed SiC whisker reinforced a-SiAlON composites’, J. Hard. Mater., 4, 121–135.

Nordberg, L.O., Ekström, T. and Xu, F.F. (1997a), ‘Simultaneously MoSi2 and SiC whiskerreinforced a-SiAlON composites’, J. Mater. Sci. Lett., 16, 917–920.

Nordberg, L.O., Shen, Z.J., Nygren, M. and Ekström, T., (1997b), ‘On the extension ofthe a-SiAlON solid solution range and anisotropic grain growth in Sm-doped a-SiAlON ceramics’, J. Eur. Ceram. Soc., 17(4), 575–580.

Oyama, Y. and Kamigaito, O., (1971), ‘Solid solubility of some oxides in Si3N4’, Jpn. J.Appl. Phys., 10, 1637–1642.

Phillips, D.C., Park, N. and Lee, R.J., (1990), ‘The impact behaviour of high performanceceramic matrix fibre composites’, Comp. Sci. and Tech., 37(1–3), 249–265.

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19.1 Introduction

Engineering ceramics based on silicon nitride are well known as low-densitymaterials with high strength and toughness. With the combination of theseproperties, silicon nitride ceramics are an ideal candidate for several structuraland functional applications. However, because of the relatively high brittlenessof ceramics there is a continual need to improve their mechanical characteristics.At the same time, intensive research is being performed to improve thethermal and electrical properties of ceramics. In general, there are two waysto improve the mechanical properties of ceramics: controlling themicrostructure, and preparation of the composite.

In connection to the microstructure–property relationship, new observationshave been performed on structural and morphological development of siliconnitride ceramics [1, 2]. It was found that the development of an interlockingmicrostructure of elongated grains is vital to ensure that this family of ceramicshas good damage tolerance. A fast (within minutes) in situ formation of atough microstructure has been observed by a so-called dynamic Ostwaldripening process that results from the rapid heating rate. In this way, throughformation of a tough interlocking microstructure (e.g. elongated b-Si3N4

grains), mechanical properties may be improved.On the other hand, physical and mechanical properties of ceramics can be

improved through nanocomposite processing [3–5]. To increase the fracturetoughness, various energy-dissipating components have been incorporatedinto ceramic matrices [4]. These secondary constituents can be introduced inwhisker, platelet, particle or fiber forms. Depending on processing route,micro/nano or nano/nano type microstructures can be synthesized, and siliconnitride–silicon carbide nanocomposites have been developed, which retainedhigh strength and good oxidation resistance up to 1400∞C [6]. Moreover,through nanocomposite processing not only the high-temperature propertiesmay be improved, but high-performance structural materials can be synthesized[7–9]. From the literature, several methods are known for porous ceramic

19Carbon-ceramic alloys

C B A L Á Z S I, Research Institute of Technical Physics andMaterials Science, Hungary

Carbon-ceramic alloys 515

preparation, such as partial sintering [8] or using sintering additives withhigh melting points to hinder the sintering process [9]. It has been reportedthat by the addition of small amounts of carbon, high-porosity silicon nitridecomposites with low shrinkage have been realized [10]. This process involvedthe formation of SiC particles at the grain boundaries, which inhibited therearrangement of Si3N4 particles during sintering, assuring high porosity. Amethod for producing porous silicon nitride with high strength and lowthermal conductivity by adding b-Si3N4 crystal seeds was presented byToriyama et al. [11, 12]. Careful control of the oxidation process of Si3N4

that contains organic binder as a source of carbon can lead to porous ceramicswith high strength [13].

As carbon nanotubes present exceptional mechanical, superior thermaland electrical properties in general, by using them as reinforcing elementsthere are high expectations for improvement of quality of nano- andmicrocomposites [14–18]. As shown from earlier measurements, throughcarbon nanotube addition a 15–37% improvement of mechanical properties(elastic modulus and strength) can be achieved in comparison to other carbon-filled samples [19].

This chapter describes the preparation and examination of ceramic matrixcomposites realized by the addition of different carbon polymorphs (carbonblack nanograins, graphite micrograins, carbon fibers and carbon nanotubes)to silicon nitride matrices. In the following sections, structural, morphologicaland mechanical characteristics of carbon-containing silicon nitride ceramicsare presented.

19.2 Carbon as fugitive additive for porous silicon

nitride processing

Several methods are known for the production of porous ceramics, such aspartial sintering [8] or using specific sintering additives [9]. An alternativemethod is to use fugitive additives such as carbon or other carbon-containingmaterials as additions, which has proven to be a successful technique for thepreparation of high-performance porous ceramics [20].

The first technological step is mixing of powders, which can be monitoredstep by step by infrared observations. Infrared spectroscopy results are presentedin Fig. 19.1. Samples from batches with carbon black and graphite additionsare characterized mainly by Si–N vibration modes (Figs 19.1(a) and (b). Inthe case of the batch with excess oxygen (Fig. 19.1(c)), due to surfaceoxidation of alpha silicon nitride powders, vibration modes of Si–O bondsappeared at 1250 cm–1 and 1100 cm–1. However, some of the vibrationmodes of the ethanol that serves as mixing agent can be recognized as OH,H2O, CH3 and C = O stretching (n) and bending (d) vibrations [21, 22].

After performing the mechanical activation of powder mixtures, rectangular


bars were obtained by dry uniaxial pressing at 220 MPa. The as-obtainedsamples were oxidized for 2 h. Oxidation at different temperatures resultedin samples with different carbon contents as presented in Fig. 19.2. Carbon

Rel

ativ

e in

ten

sity

c

b

a

Si–N

Si–O

CO2

4000 3000 2000 1000 0Wavenumber (cm–1)

19.1 FTIR spectra of milling products: (a) resulting mixture withcarbon black after milling; (b) resulting mixture with graphite aftermilling; (c) mixture after milling, batch with oxidized starting powder.

GraphiteCarbon black

Oxidizedstarting powder

12

10

8

6

4

2

0

Car

bo

n c

on

ten

t (w

t%)

400 450 500 550 600 650 700 750 800 850Temperature (∞C)

19.2 Samples from batches with carbon black, graphite and oxidizedstarting powder with different carbon content after oxidation inatmosphere. Average values of four samples are presented.


removal was very fast in the case of the batch with the oxidized surface. Thebehavior of the carbon black and graphite added silicon nitride matrix wasfound to be significantly different regarding the oxidation process. In thetemperature range from 450∞C up to 600∞C the carbon black content decreasedfrom 12 wt% to 0.4 wt%. At this stage the graphite content was around 10wt%. Above this temperature the graphite content also had a tendency todecrease with increasing temperature. Continuing the oxidation process above800∞C, according to weight losses we obtained almost a carbon-free structure.During the oxidation process carbon is released as CO or CO2 gas while thesurface of a-Si3N4 in the compacts may also be oxidized, but we have notobserved Si–O bonds evolving in the infrared spectra of compacts.

A de-sintering process (i.e. decreasing of density) was observed duringthe two-step HIP sintering process, as shown in Fig. 19.3. The starting pointsof arrows represent the first sintering step without any pressure applied. Theends (tips) of arrows represent the second sintering step with 2 MPa pressureapplied. The de-sintering process was observed for the reference sample andfor batches with carbon black (samples containing ~0.1 wt% carbon afteroxidation) and graphite (e.g. samples with no graphite content after oxidationat 800∞C) as well. A very similar de-sintering phenomenon was described byHwang et al. in the case of gas-pressure sintering [23]. The decrease ofdensity was ascribed to chemical dissolution of nitrogen into oxynitride

With carbon black

With graphite

Referencesilicon nitride

2.0 2.5 3.0 3.5Apparent density (g/cm3)

300

250

200

150

100

50

Mo

du

lus

of

elas

tici

ty (

GP

a)

19.3 Effect of de-sintering observed during two-step sinteringprocess (without and with pressure applied) on reference sample,two samples with ~0.1 wt% carbon content after oxidation at 800∞Cfrom batch with carbon black, and one sample from batch withgraphite and with no graphite content after oxidation at 800∞C.


melts assisted by high pressure and presence of boron nitride phases. Aninteresting remark should be added to the de-sintering observation presentedin Fig. 19.3. In the case of the reference sample we found a similar decreaseof density as in the batches with carbon black and graphite, but the effect onmodulus of elasticity is reversed. A comprehensive view on the carbon contenteffect on strength can be seen in Fig. 19.4. Samples with added graphiteshow higher values of strength as compared to carbon black and oxidizedsamples.

19.3 Comparison of silicon nitrides with carbon

additions prepared by hot isostatic pressing

and pressureless sintering

In addition to dense monolithic ceramics, porous silicon nitrides are gainingmore importance in technological applications [24]. Some porous siliconnitrides with high specific surface area have already been applied as catalysissupports, hot gas filters and biomaterials [25]. There is an emerging tendencyto facilitate silicon nitride as biomaterial, because of specific mechanicalproperties that are important for medical applications [25]. Moreover, in arecent study it was shown that silicon nitride is a non-toxic, biocompatibleceramic which has the ability to propagate human bone cells in vitro [25].Bioglass and silicon nitride composites have already been realized to combine

BS4BS3With graphite

With carbon blackOxidized starting powder

Ben

din

g s

tren

gth

(M

Pa)

300

250

200

150

100

50

00 1 8 12 0 10 12 0 6 10 12 15

Carbon content (wt%)

19.4 Relation between carbon content and four-point (BS4) and three-point bending strength (BS3).


their advantageous properties, such as excellent bioactivity, biocompatibilityand favorable mechanical characteristics, particularly the required fracturetoughness and wear resistance of components that are prerequisites for clinicalapplications [26, 27]. Porous silicon nitride rods inserted into rabbit femurshave been found to promote bone ingrowth. In this way a possible newapplication of silicon nitride is suggested, namely to promote an ideal candidatefor bone-substituting implants.

Details of the composition of the starting powder mixtures and preparationcan be seen in Table 19.1. Si3N4 (Ube, SN-ESP), Al2O3 (Alcoa, A16) andY2O3 (H. C. Starck, grade C) were used as starting powders. For compositeprocessing in addition to batches, carbon black (Taurus Carbon black, N330,average particle size ~50–100 nm) and graphite (Aldrich, synthetic, averageparticle size 1–2 mm) were added. The powder mixtures were milled inethanol in a planetary-type alumina ball mill for 3 h. It was found fromweight measurements that each batch contained approximately 1 wt% aluminaas contamination from balls and jars. The batches were dried and sieved.Green samples were obtained by dry uniaxial pressing at 220 MPa. Beforesintering an oxidation was applied at very low heating rates up to 400∞C, toeliminate the compaction agent, polyethyleneglycol (PEG), from samples.

Hot isostatic pressing (HIP) was performed at 1700∞C under a pressure of20 bar in high-purity nitrogen by a two-step sinter-HIP method using BNembedding powder (Ceramic and Composites Laboratory, Budapest). Theheating rate did not exceed 25∞C/min. Pressureless sintering (PLS) wasperformed in a graphite furnace (Thermal Technology GmbH) at 1700∞C for2 hours under 1 bar flowing nitrogen (Ceramic Research Unit, Limerick).The dimensions of the as-sintered specimens were 3.5 ¥ 5 ¥ 50 mm. Aftersintering, the weight change of the samples was determined. All surfaces ofthe samples were finely ground on a diamond wheel, and the edges werechamfered. The density of the sintered materials was measured by theArchimedes method. Phase compositions were determined by Philips PW1050 diffractometer. The morphology of the solid products was studied byfield emission scanning electron microscope, LEO 1540 XB. For HIPedsamples the elastic modulus, four-point and three-point bending strengthswere determined by a bending test with spans of 40 and 20 mm. Hardnessmeasurements were performed in a micro Vickers Model LL, Tukor Tester,by applying 1 kg load.

Fracture surfaces of samples prepared by pressureless sintering are shownin Fig. 19.5. As can be observed, the preparation conditions assured amicrostructure with pores of 2–5 mm in dimensions even in the case of thereference sample. Using carbon black or graphite addition, the pore sizecould be increased (Fig. 19.5(a), (b). At the same time, by increasing thecarbon addition the morphology (and microstructural aspects) could be changed.Large silicon carbide granules can be observed in the center of Fig. 19.5(b),

Ceram

ic matrix com

posites520

Table 19.1 Starting compositions and preparation conditions of sintered samples

Treatment* Sample Starting powder Carbon Ball milling Sintering conditions(wt%) black/graphite (in ethanol) ——————————————————

———————————————— addition time Temp. Holding PressureSi3N4 Al2O3 Y2O3 (wt%)† (∞C) time (bar)

PLS PLS-1 90 4 6 — 3 h 1700 2 h 1PLS-3 90 4 6 1CB 3 h 1700 2 h 1PLS-5 90 4 6 10CB 3 h 1700 2 h 1PLS-7 90 4 6 1G 3 h 1700 2 h 1PLS-9 90 4 6 10G 3 h 1700 2 h 1

HIP HIP-1 90 4 6 – 3 h 1700 — 20HIP-2 90 4 6 1CB 3 h 1700 — 20HIP-3 90 4 6 10CB 3 h 1700 — 20HIP-4 90 4 6 1G 3 h 1700 — 20HIP-5 90 4 6 10G 3 h 1700 — 20

*PLS pressureless sintering, HIP hot isostatic pressing.†CB carbon black, G. graphite.


10mmMag = 1.00 K X

EHT = 1.00 kVWD = 8 mm

Signal A = SE2Photo No. = 7605

Date: 15 Jun 2004Time: 9:28

(a)

(b)

20mmMag = 1.00 K X

EHT = 2.00 kVWD = 7 mm

Signal A = SE2Photo No. = 7620

Date: 15 Jun 2004Time: 10:49

19.5 (a) Sample PLS-5, 10% CB with large pores; (b) sample PLS-9,10% G with pores and large platelets of SiC.


and in this connection we found that the structural characteristics had beenchanged, according to Fig. 19.6.

As resulting from XRD measurements, the lines of a-sialon (PDF-JCPDS42-0251), together with a-Si3N4 (PDF-JCPDS 41-0360), mellilite (Y2Si3N4O3,PDF-JCPDS 30-1460) and small additions of b-Si3N4 (PDF-JCPDS 33-1160),can be observed (sample PLS-1). After 1% carbon addition (Fig. 19.6, samplesPLS-3 and PLS-7) the same main structural lines could be observed as in thereference sample (PLS-1). After we increased the carbon to 10% (PLS-5 andPLS-9) the a-sialon and melilite lines disappeared. In this case the mainconstituents of the microstructures are a-Si3N4, b-Si3N4 and SiC (PDF-

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1700∞C, 1 bar, PLS

PLS-1

PLS-7, 1%G

PLS-9, 10%G

PLS-3, 1%CB

PLS-5, 10%CB

0 10 20 30 40 50 60 70 802q, Cuka

1700∞C, 20 bar, HIP

0 10 20 30 40 50 60 70 802q, Cuka

HIP-1

HIP-4, 1%G

HIP-5 10%G

HIP-2, 1%CB

HIP-3, 10%CB

19.6 X-ray measurements for samples resulting from PLS and HIP.


JCPDS 31-1231). For samples prepared by the HIP process we found onlya-Si3N4 and b-Si3N4 as constituents for all of the structures (Fig. 19.6).

A comparison of three-point bending strengths is presented in Fig. 19.7.By increasing the carbon content an increase of porosity and a decrease ofbending strength can be observed for PLSed (open symbols) and HIPedsamples (filled symbols). Contrary to observations made in hot pressing[28], we found an increase of hardness with an increase in the carbon content(Table 19.2). The scatter (related to density and strength as well) of PLSedsamples is smaller than for HIPed samples. In the small density ranges (to2.7 g/cm3), by changing from PLS to HIP samples with higher densities andhigher strengths have been obtained. At higher density levels the changefrom PLS to HIP is accompanied by lower densities, but with higher strengths.

19.7 Three-point bending strength of reference sample andcomposites as function of apparent density. At each point theaverage value of five samples is shown.

Table 19.2 Hardness values for samplesprepared by PLS

Samples HV (GPa)

PLS-5, 10%CB 5.6 ± 0.9PLS-3, 1%CB 15.9 ± 0.7PLS-9, 10%G 21.2 ± 0.6PLS-7, 1%G 18.0 ± 0.8PLS-1, 0% 11.0 ± 0.7

0 0.5 1.0 1.5 2.0 2.5 3.0 3.5Apparent density (g/cm3)

600

500

400

300

200

100

0

Th

ree-

po

int

ben

din

g s

tren

gth

(M

Pa)

0% 1700∞C 1 bar0% 1700∞C 20 bar1%CB 1700∞C 1 bar1%CB 1700∞C 20 bar1%G 1700∞C 1 bar1%G 1700∞C 20 bar10%CB 1700∞C 1 bar10%CB 1700∞C 20 bar10%G 1700∞C 1 bar10%G 1700∞C 20 bar


19.4 In situ processing of Si3N4/SiC composites by

carbon addition

Several research groups reported a new type of silicon nitride–silicon carbidenanocomposite with improved high-temperature strength and fracture toughness[29, 30]. Niihara and coworkers [31] were the first research group to studythe fabrication of Si3N4/SiC nanocomposites in detail. Their production routeconsisted of hot pressing an amorphous SiCN powder obtained via a chemicalvapor deposition process. These composites consist of a submicron-sizedSi3N4 matrix that contains a dispersion of nano-sized and globular-shapedSiC particles. The particles are located both at Si3N4 grain boundaries andwithin Si3N4 grains. Si3N4/SiC nanocomposites exhibit excellent mechanicalproperties at both ambient and elevated temperatures. Different strategies forthe synthesis of silicon carbide–silicon nitride composites from pre-ceramicpolymers, polymer-derived powders or other novel methods have been presented[32, 33]. The polymeric precursor route to synthesize ceramic compositesoffers processing at low temperature and the possibility of optimizing thenew composite materials production by tailoring the composition and molecularstructure of the polymers. Recently, a low-cost silicon carbide–silicon nitridenanocomposite processing route has been reported [34] in which case, thebulk silicon nitride-based nanocomposite is formed by the carbothermalreduction of SiO2 by carbon in the Y2O3–SiO2 system at the sinteringtemperature, with SiC nanoparticles as the result. Although the mechanicalproperties of as-prepared samples should be further optimized, this processseems to be a prospective choice for silicon nitride–silicon carbidenanocomposite production.

19.4.1 Size effects in micro- and nano-carbon addedC/Si3N4 composites prepared by hot pressing

The powder mixtures (Table 19.3) with carbon black additions were milledin ethanol in a planetary-type alumina ball mill (Fritsch GmBH). For graphite-added mixtures a highly efficient attritor mill (Union Process, type 01-HD/HDDM) was employed. This apparatus allowed a higher rotation speed anda contamination-free mixing process, because of silicon nitride parts (tank,arm, balls). Samples for hot pressing (HP, CENTORR Vacuum Industries)were prepared as follows. After 10–4 torr vacuum, at 20∞C, nitrogen gas wasintroduced at 1 atm. The heating rate was 20∞C/min up to 1800∞C. At 1000∞Ca uniaxial pressure of 2 MPa was applied and kept at this pressure until1800∞C. At 1800∞C, the pressure was applied for 2 h. The samples werecooled together with the furnace (cooling rate was 30∞C/min).

In Fig. 19.8(a) and Fig. 19.8(c) the fine microstructures of hot-pressedreference samples resulting from ball and attritor milling and sintering are

Carbon-ceram

ic alloys525

Table 19.3 Preparation conditions and composition of starting powder mixtures

Batch Composition (wt%) Carbon black/graphite Ball milling Attritor milling Hot pressing 1800∞C,—————————————— added to batch (ethanol) at (ethanol) at 2 h nitrogen atm.Si3N4 Al2O3 Y2O3 (wt%) 360 rpm 600 rpm (MPa)

CB0% 90 4 6 — 3 h — 2CB1% 90 4 6 1 3 h — 2CB10% 90 4 6 10 3 h — 2

G0% 90 4 6 — — 3 h 2G1% 90 4 6 1 — 3 h 2G10% 90 4 6 10 — 3 h 2


shown. Both microstructures consisted of submicron grains, though in thecase of the reference sample (CB0%) small hexagonal-shaped b-Si3N4 grainshad appeared. Indeed, as presented in Fig. 19.9, in the microstructure ofsample CB0% a considerable amount of b-Si3N4 lines appeared. A porousmicrostructure and an enhanced grain (b-Si3N4) growing process can beobserved in the CB10% scanning micrograph (Fig. 19.8(b). Graphite plateletswere included in the microstructure even after sintering, as shown in Fig.19.8(d), sample G10%. By adding 1 wt% carbon black the microstructureremained the same as for the reference sample: a-Si3N4 (PDF-JCPDS 41-0360) and b-Si3N4 (PDF-JCPDS 33-1160) were the main phases (Fig. 19.9,sample CB1%), but with an increase in carbon black (CB10%) SiC lines hadappeared (PDF-JCPDS 31-1231). This observation was confirmed by promptgamma activation analysis (PGAA) as well. This method is based on thedetection of prompt gamma rays originating from the (n, g) reaction, andgives average elemental composition of the total volume of the sample [28,35]. The compositions of some sintered samples are presented in Table 19.4.

(a) (b)

(c) (d)

19.8 Micrographs of fracture surfaces: (a) reference sample CB0%,(b) sample CB10%; (c) reference sample G0%, (d) sample G10%.Bar: 1 mm for all micrographs.


By comparing the Si/N mass ratios of the monolithic reference sample andthe composite sample, the effect of carbon content on the complex sinteringprocess can be determined.

The results show that at low pressure in the presence of 10 wt% carbonthe Si/N mass ratio had considerably increased (to 1.91 ± 0.06 for CB10%)

Rel

ativ

e in

ten

sity

Rel

ativ

e in

ten

sity

SiC

c

b

a

0 10 20 30 40 50 60 70 802q, Cuka

c

b

a

0 10 20 30 40 50 60 70 802q, Cuka

19.9 X-ray diffractograms of sintered samples. Top panel: (a)reference sample CB0%, (b) sample CB1%, (c) sample CB10%.Bottom panel: (a) reference sample G0%, (b) sample G1%, (c) sampleG10%.


compared to the reference sample (from 1.49 ± 0.04 for CB0%). The Si/Nfraction for other samples was close to the theoretical value for an Si3/N4

ratio equal to 1.5. In accordance with X-ray measurements, which show theformation of SiC at low pressure and 10 wt% carbon content, the increase ofSi/N ratio revealed a nitrogen exhaust from the structure according to thefollowing reaction taking place around 1435ºC [28]: 3C(s) + Si3N4 (s) Æ3SiC(s) + 2N2(g). In the case of graphite additions, a-Si3N4, b-Si3N4 and b-Y2Si2O7 (PDF 21-1454 JCPDS) lines can be observed (Fig. 19.9, samplesG1% and G10%). The SiC phase was not formed in graphite composites.Graphite addition assured an increase of hardness in comparison to samplewith carbon black added (Table 19.5). Addition of graphite resulted in moredensified samples than for samples containing carbon black (Fig. 19.10). Allmodulus values have been found to be higher for graphite than for carbon

Table 19.4 Compositions of sintered (at 2 MPa) reference sample CB0%and CB10%

Reference sample CB0%Elements Wt% Rel. unc.%*

H 0.0048 12.9B 0.0093 1.0C — —N 36.6853 2.1F 1.0272 29.1Al 2.3548 2.5Si 54.7314 2.5S 0.0240 33.1Cl 0.0096 3.6Y 5.1402 1.9Si/N = 1.49 ± 0.04

Sample CB10%Elements Wt% Rel. unc.%*

H 0.0263 5.9B 0.0178 1.0C 12.952 10.4N 24.192 2.3F 9.4493 8.4Al 2.7066 3.0Si 46.423 2.5S — —Cl 0.0014 24.1Y 4.2161 1.9Si/N = 1.91 ± 0.06

*Because of high relative uncertainty oxygen was not taken intoconsideration.


black added composites at the same carbon level. As regards bending strengths,although the reference sample has higher strength, for graphite addedcomposites a higher strength was obtained than for carbon black addedsamples (Fig. 19.10). We found an increase in strength for G1% in comparisonto reference G0%. It seems that, because of the high specific surface area of

Table 19.5 Hardness of composites

Batch Hardness, HV (GPa) Batch Hardness, HV (GPa)

CB0% 13.07 ± 0.6 G0% 16.27 ± 0.6CB1% 9.22 ± 0.7 G1% 14.72 ± 1.2CB10% 1.21 ± 0.5 G10% 8.94 ± 2.1

1.8 2.0 2.2 2.4 2.6 2.8 3.0 3.2Apparent density (g/cm3)

300

250

200

150

100

50

0

Mo

du

lus

(GP

a)

2MPa + 0% Graphite2MPa + 1% Graphite2MPa + 10% Graphite2MPa + 0% Carbon black2MPa + 1% Carbon black2MPa + 10% Carbon black

1.8 2.0 2.2 2.4 2.6 2.8 3.0 3.2Apparent density (g/cm3)

600

500

400

300

200

100

0

Th

ree-

po

int

ben

din

g s

tren

gth

(M

Pa)

19.10 Modulus of elasticity and three-point bending strength asfunctions of density.


carbon black nanograins, agglomeration occurs that causes a more porousmicrostructure after sintering, with lower mechanical properties.

19.5 Silicon nitride ceramics reinforced with carbon

fibers and carbon nanotubes

Carbon nanotubes (CNTs) present exceptional mechanical, superior thermaland electrical properties; therefore, in general, there are high expectationsfor improvement of quality of carbon nanotube nano- and microcomposites[14–18]. Despite this, only modest improvements have been reported inrelation to electrical or mechanical properties of carbon nanotube siliconcarbide-, polymer- or metal oxide-matrix composites [36–39]. These worksemphasized that research concerning CNT-reinforced ceramic matrixcomposites should deal with increase of dispersion grade of CNTs in theceramic matrix, assurance of good CNT/matrix interconnection, and high-temperature protection of CNTs. One of the possible approaches is to modifythe shape factor of CNTs [40]. It was shown that an optimization of the stressdistribution and the reinforcement of mechanical bonding can be achievedby using non-circular or hollow carbon fibers in an epoxy matrix [41]. In arecent study, high-temperature extrusion has been used to align the CNTs inbulk metal oxide nanocomposites, which could lead to an increase of electricalconductivity [42]. Many efforts have been made to modify the surface propertiesof carbon nanotubes. Using an acidic treatment in a cc.H2SO4/cc.HNO3

mixture results in nanotubes covered by carboxyl groups at their ends and/or their shells. These groups can easily be converted into carbonyl chloridegroups simply by reacting them with SOCl2. The resulting material is veryreactive towards amines, so by choosing the appropriate reactant any kind offunctional group can be generated [43]. A simple mechano-chemicalfunctionalization of the CNT surface [44] or application of different coatings[45] may be prospective choices to enhance nanotube–matrix bonding.

An alternative way to enhance the interfacial strength is to use a physicalactivation of CNT by spark plasma sintering (SPS) processing. As alreadyshown, good interface contact was achieved in alumina–CNT composites bySPS [1]. This good contact may be attributed to a surface activation, whichis assumed to be due to the external electrical field application in SPS. Whilethe surface activation was shown in different materials, no systematic studyhas been performed. The surface activation and enhanced densification infield sintering is most noticeable at lower temperatures when the dischargeeffects are operational. The highest temperatures achieved in the necks providethe highest diffusion rates and thus enhanced matter transport towards theneck area. This is the area in which most of the matter transport is requiredfor sintering. Therefore, field application intensifies the sintering rate. Beyondthe electrical discharge stage, the sintering parameters in SPS have a similar


effect as in conventional pressure sintering. Diffusion processes and plasticflow are the main contributors to the densification. Diffusion studies indicatethe extensive contribution of an applied electrical current (e.g., electromigration)in all sintering stages. In the initial sintering stages, the local dischargeprocesses, which disrupt the surface layers, induce various types of defectsthat enhance surface and grain boundary diffusion. This way, both densificationand grain growth are accelerated. Basic studies showed that field applicationis likely to enhance grain boundary velocity due to impurity elimination.

The manufacture of a ceramic composite generally requires hightemperatures. Destruction of carbon nanotubes using a hot-press techniquewas reported by Flahaut et al. [39]. Therefore, the high-temperature degradationprocess of CNTs has to be further optimized to achieve proper protection ofCNTs in high-temperature processes.

Concerning carbon nanotube-reinforced silicon nitride matrices, only afew reports have so far been published [19]. In this case, hot isostatic pressinghas been used for composite processing. The carbon nanotubes remained inthe microstructure only under low pressures (2 MPa); they connect the siliconnitride grains and produce a 15–37% improvement of the mechanical propertiesas compared with other carbon-filled samples (Fig. 19.11). Increase of pressure

Table 19.6 Mechanical properties (elastic modulus (E), shear modulus (G),Poisson’s ratio (n), hardness (HV), and fracture toughness (KIC) of compositessintered by SPS

Samples/ r (g/cm3) E (GPa) G (GPa) n HV KIC

SPS (GPa) (MPa.m1/2)

Reference 3.23 326.21 130.18 0.25 20.1 ± 0.9 5.2CNT/Si3N4 3.17 285.73 115.10 0.24 16.6 ± 0.4 5.3

4 point bending strength3 point bending strength

Apparent density (g/cm3)2,315 2,415 2,42 2,467

350

300

250

200

150

100

50

0

Ben

din

g s

tren

gth

(M

Pa)

629644

645

684

19.11 Carbon nanotube-ceramic composites (684) showing 37% increasein bending strength compared with other carbon-filled samples (629carbon fiber, 644 carbon black, 645 graphite), from Ref. 19.


and sintering time resulted in the disappearance of the nanotubes. SPS sinteringhas a proven ability to sinter powders at lower temperatures than other sinteringmethods. The CNTs are located mainly in the intergranular places and presentgood adherence to silicon nitride grains (Fig. 19.12(b)). Pullouts can also benoted on the fracture surface, though as can be seen, the CNTs are connected

19.12 Scanning electron micrographs of fracture surfaces of SPSsamples: (a) reference sample sintered at 1500ºC for 3 min at 50MPa; (b) sample with MWNTs dispersed and located between grains,sintered at 1500ºC for 5 min at 100 MPa.

(a)

(b)


to each other and can be found in groups and nano- or micron–sized islandseven after sintering. Therefore, samples with added CNT are characterizedby greater porosity than reference samples, and the local reinforcementsnoted as pullouts are not enough for significant global strengthening. As canbe observed in Table 19.6, the reference samples with higher densities showhigher modulus, hardness and toughness. Keeping in mind the complexity ofthe consolidation process of silicon nitride, more precise control is needed toquantify the contribution of carbon nanotubes to the end properties ofcomposites.


Based on the new types of ceramic matrix discussed above, composites withimproved properties can be realized by carbon addition. Silicon nitrides withcontrolled porosity and in situ silicon carbide/silicon nitride composites wereobtained with the help of carbon additions and by tailoring the complexreaction paths within sintering processes. Carbon nanotubes have beensuccessfully applied as ceramic matrix reinforcements. The manufacturingprocesses have been optimized in order to thoroughly disperse the carbonnanotubes in the matrix, to assure good nanotube–silicon nitride contact, andto keep intact the nanotubes during high-temperature processing. An idealprocessing–microstructure–property relationship was developed anddemonstrated.

19.7 References

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2. Shen, Z., Zhao, Z., Peng, H., Nygren, M., Formation of tough interlockingmicrostructures in silicon nitride ceramics by dynamic ripening, Nature, 417, 16May 2002, 266–269.

3. Dobedoe, R.S., West, G.D., Lewis, M.H., Spark plasma sintering of ceramics, Bull.Eur. Ceram. Soc., 1, 2003, 19–24.

4. Sternitzke, M., Structural ceramic nanocomposites, J. Eur. Ceram. Soc., 17, 1997,1061–1082.

5. Derby, B., Ceramic nanocomposites: mechanical properties, Cur. Op. in Solid StateMat. Sci., 3, 1998, 490–495.

6. Niihara, K., New design concept of structural ceramics – ceramic nanocomposites,J. Jpn. Ceram. Soc., 99(10), 1991, 974–982.

7. Rendtel, A., Hübner, H., Hermann, M., Schubert, C., Silicon nitride/silicon carbidenanocomposite materials: II, Hot strength, creep and oxidation resistance, J. Am.Ceram. Soc., 81(5), 1998, 1109–1120.

8. Arató, P., Besenyei, E., Kele, A., Wéber, F., Mechanical properties in the initial stageof sintering, J. Mat. Sci., 30, 1995, 1863–1871.


9. Yang, J.-F., Deng, Z.-Y., Ohji, T., Fabrication and characterisation of porous siliconnitride ceramics using Yb2O3 as sintering additive, J. Eur. Ceram. Soc., 23, 2003,371–378.

10. Yang, J.-F., Yhang, G.-J., Ohji, T., Fabrication of low-shrinkage, porous siliconnitride ceramics by addition of a small amount of carbon, J. Am. Ceram. Soc., 84(7),2001, 1639–1641.

11. Toriyama, M., Hirao, K., Brito, M.E., Kanzaki, S., Shigegaki, Y., US Patent no.5,935,888.

12. Hirao, K., Brito, M.E., Toriyama, M., Kanzaki, S., Imamura, H., Hirai, T., Shigegaki,Y., US Patent no. 5,968,426.

13. Kawai, C., Matsuura, T., Yamakawa, A., US Patent no. 5,780,374.14. Rochie, S., Carbon nanotubes: exceptional mechanical and electrical properties,

Ann. Chim. Sci. Mat., 25, 2000, 529–532.15. Thostenson, E.T., Ren, Z., Chou, T.W., Advances in the science and technology of

carbon nanotubes and their composites: a review, Comp. Sci. and Techn., 61, 2001,1899–1912.

16. Lau, K.T., Hui, D., The revolutionary creation of new advanced materials – carbonnanotube composites, Composites: Part B, 33, 2002, 263–277.

17. Lourie, O., Wagner, H.D., Evidence of stress transfer and formation of fractureclusters in carbon nanotube-based composites, Comp. Sci. and Techn., 59, 1999,975–977.

18. Lau, K.T., Hui, D., Effectiveness of using carbon nanotubes as nano-reinforcementsfor advanced composite structures, Carbon, 40, 2002, 1597–1617.

19. Balázsi, Cs., Kónya, Z., Wéber, F., Biró, L.P., Arató, P., Preparation and characterizationof carbon nanotube reinforced silicon nitride composites, Mat. Sci. Eng. C, 23(6–8),2003, 1133–1137.

20. Balázsi, Cs., Cinar, F.S., Addemir, O., Wéber, F., Arató, P., Manufacture and examinationof C/Si3N4 nanocomposites, J. Eur. Ceram. Soc., 24 (12), 2004, 3287–3294.

21. Balázsi, Cs., Wéber, F., Arató, P., Investigation of C/Si3N4 nanocomposites, Mat.-wiss. u. Werkstofftech, 34, 2003, 332–337.

22. Nakamoto, K., Infrared Spectra of Inorganic and Coordination Compounds, Wiley-Interscience, New York, 1963.

23. Hwang, S.-L., Becher, P.E., Lin, H.T., Desintering process in the gas pressure sinteringof silicon nitride, J. Am. Ceram. Soc., 80(2), 1997, 329–335.

24. Diaz, A., Hampshire, S., Characterisation of porous silicon nitride materials producedwith starch, J. Eur. Ceram. Soc., 24, 2004, 413–419.

25. Kue, R., Sohrabi, A., Nagle, D., Frondoza, C., Hungerford, D., Enhanced proliferationand osteocalcin production by human osteoblast-like MG63 cells on silicon nitrideceramic discs, Biomaterials, 20, 1999, 1195–1201.

26. Amaral, M., Lopes, M.A., Silva, R.F., Santos, J.D., Densification route and mechanicalproperties of Si3N4–bioglass biocomposites, Biomaterials, 23, 2002, 857–862.

27. Guedes e Silva, C.C., Higa, O.Y., Bressiani, J.C., Cytotoxic evaluation of siliconnitride-based ceramics, Mat. Sci. Eng. C, 24, 2004, 643–646.

28. Balázsi, Cs. Cinar, F.S., Kasztovszky, Zs., Cura, M.E., Yesilcubuk, A., Wéber, F.,Investigation of hot pressed C/Si3N4 nanocomposites, Silicates Industriels, 69 (7–8),2004, 293–298.

29. Pezzotti, G., Tanaka, I., Okamoto, T., Si3N4/SiC-whisker composites without sinteringaids: III, High-temperature behavior, J. Am. Ceram. Soc., 74(2), 1991, 326–332.

30. Goto, Y., Tsuge, A., Mechanical properties of unidirectionally oriented SiC-whisker-


reinforced Si3N4 fabricated by extrusion and hot-pressing, J. Am. Ceram. Soc., 76(6),1993, 1420–1424.

31. Niihara, K., Suganuma, K., Nakahira, A., Izaki, K., Interfaces in Si3N4–SiCnanocomposite, J. Mater. Sci. Lett., 9, 1990, 598–599.

32. Sajgalik, P., Hnatko, M., Lofaj, F., Hvizdos, P., Dusza, J., Warbichler, P., Hofer, F.,Riedel, R., Lecomte, E., Hoffmann, M.J., SiC/Si3N4 nano/micro composite – processing,RT and HT mechanical properties, J. Eur. Ceram. Soc., 20, 2000, 453–462.

33. Poorteman, M., Descamps, P., Cambier, F., Plisnier, M., Canonne, V., Descamps,J.C., Silicon nitride/silicon carbide nanocomposite obtained by nitridation of SiC:fabrication and high temperature mechanical properties, J. Eur. Ceram. Soc., 23,2003, 2361–2366.

34. Hnatko, M., Galusek, D., Sajgalik, P., Low-cost preparation of Si3N4–SiC micro/nano composites by in-situ carbothermal reduction of silica in silicon nitride matrix,J. Eur. Ceram. Soc., (2), 24 2004, 189–195.

35. Révay, Zs., Molnár, G.L., Standardisation of the prompt gamma activation analysismethod, Radiochim. Acta, 91, 2003, 361–369.

36. Laurent, Ch., Peigney, A., Dumortier, O., Rousset, A., Carbon nanotubes–Fe–aluminananocomposites. Part II: Microstructure and mechanical properties, J. Eur. Ceram.Soc., 18, 1998, 2005–2013.

37. Laurent, Ch., Peigney, A., Flahaut, E., Rousset, A., Synthesis of carbon nanotubes–Fe–Al2O3 powders. Influence of the characteristics of the starting Al1.8Fe0.2O3

oxide solid solution, Mat. Res. Bull., 35, 2000, 661–663.38. Peigney, A., Laurent, Ch., Flahaut, E., Rousset, A., Carbon nanotubes in novel

ceramic matrix composites, Ceram. Int., 26, 2000, 677–683.39. Flahaut, E., Peigney, A., Laurent, Ch., Marliere, Ch., Chastel, F., Rousset, A., Carbon

nanotube–metal-oxide nanocomposites: microstructure, electrical conductivity andmechanical properties, Acta Mater., 48, 2000, 3803–3812.

40. Park, S.J., Seo. M.K., Shim, H.B., Effect of fiber shapes on physical characteristicsof non-circular carbon fiber-reinforced composites, Mat. Sci. Eng. A, 352, 2003, 34–39.

41. Wagner, H.D., Nanotube–polymer adhesion: a mechanics approach, Chem. Phys.Lett., 361, 2002, 57–61.

42. Peigney, A., Flahaut, E., Laurent, Ch., Chastel, F., Rousset, A., Aligned carbonnanotubes in ceramics–matrix nanocomposites prepared by high-temperature extrusion,Chem. Phys. Lett., 352, 2002, 20–25.

43. Kiricsi, I., Kónya, Z., Niesz, K., Koós, A.A., Biró, L.P., Synthesis procedures forproduction of carbon nanotube junctions, Proceedings of SPIE – The InternationalSociety for Optical Engineering, 5118, 2003, 280–287.

44. Kónya, Z., Vesselényi, I., Niesz, K., Kukovecz, A., Demortier, A., Fonseca, A.,Delhalle, J., Mekhalif, Z., Nagy, J.B., Koós, A.A., Osváth, Z., Kocsonya, A., Biró,L.P., Kiricsi, I., Large scale production of short functionalized carbon nanotubes,Chem. Phys. Lett., 360, 2002, 429–435.

45. Hernadi, K., Ljubovici, E., Seo, J.W., Forró, L., Synthesis of MWNT-based compositematerials with inorganic coating, Acta. Mater., 51, 2003, 1447–1452.

536

20.1 Introduction

Despite the excellent combination of thermal mechanical properties of Si3N4

and SiC ceramics, the wide application of these materials is still hamperedby their relatively low fracture toughness (Lange, 1973; Jack, 1976; Swain,1994; Kleebe et al., 1999). Improvements in fracture toughness of Si3N4 andSiC ceramics are mainly achieved by crack deflection and bridging mechanisms(Fairbanks et al., 1987; Bennison and Lawn, 1989; Kelly et al., 1991; Chenand Engineer, 1999). To enhance crack deflection toughening, high aspectratio acicular grains and a weak interface between the grains and grainboundary phases are desirable; to enhance crack bridging toughening, highaspect ratio grains with large diameters are required. However, weak interfacesand subsequently the debonding of large elongated grains result in thedevelopment of fracture origins and loss of strength of the material. Therefore,to develop a superior material with both high toughness and strength,optimization of the microstructure is essential. Accordingly, toughness andstrength have been areas of intensive research over the past two decades. Theoutcomes have been quite impressive: materials with much improved fracturetoughness and strength, exceeding 12 MPa m1/2 and 1 GPa respectively, havebeen reported (Hirao et al., 1995; Ohji et al., 1995; Becher et al., 1998).

Although it has long been recognized that microstructure has a significanteffect on the erosion performance of ceramic materials, the role ofmicrostructure in the erosion process is not understood in detail. Most of theprevious studies on the effect of microstructure on erosion response usedalumina ceramics as the model material. However, alumina ceramics, ingeneral, contain an equiaxed grain morphology and high residual stresses atthe grain boundaries and are, therefore, not representative of other engineeringceramics, especially the self-reinforced ceramics. In addition, many of theprevious studies were carried out on commercial materials in which there islittle scope for independent control of microstructures. Thus in the currentresearch, the role of microstructure in the erosion process has been studied

20Silicon nitride and silicon carbide-based

ceramics

Y Z H A N G, New York University College of Dentistry, USA

Silicon nitride and silicon carbide-based ceramics 537

using a variety of nitride and carbide ceramics with different microstructures,including several in-house prepared sialon ceramics. Sialon ceramics aregenetically associated with Si3N4, but have greater phase complexity andmore degrees of freedom for tailoring of microstructures. This offers an idealsystem for studying the role of microstructure on erosion of advanced ceramics.

20.2 Material selection

20.2.1 Sialons

a-Sialons are isostructural with a-Si3N4 and can be described by the formula

M +xv Si12–(m+n)Alm+nOnN16–n (20.1)

where M is the compensating cation with a valency n, and m and n aresubstitution numbers referring to m (Al–N) and n (Al–O) bonds replacing(m + n) (Si–N) bonds in each unit cell (Cao and Metselaar, 1991). x is relatedto m according to the relationship x = m/n and has a value £2 (Hampshire etal., 1978). Two Ca a-sialon compositions, namely CA1005 and CA2613,with nominal x (= m/2 = n) values of 0.5 and 1.3, respectively, were selectedfor this study. Composition CA1005 lay inside the single-phase a-sialonforming region of the Ca a-sialon-phase diagram, while composition CA2613was situated just outside the single-phase region on the Al-rich side. Detailsof the compositional design and powder preparation procedures of thesematerials are described in a previous paper (Zhang and Cheng, 2003). In thisstudy, both compositions were first pressureless-sintered (PLS-ed) at 1800∞Cfor 3 h followed by hot pressing (HP) at 1700∞C for 1 h.

Typical microstructures of polished sections of the Ca a-sialon materialsimaged in a SEM using the secondary electron mode are shown in Fig. 20.1.Sample CA1005 contained almost equiaxed a-sialon grains with a fractionof the grains being slightly elongated (Fig. 20.1(a)). The average diametersof the a-sialon grains were 0.52 mm. Image analysis revealed that the volumefraction of grain boundary glass was ~2%. No secondary crystalline phasewas detected in these materials according to the XRD analysis. XRD analysisshowed that material CA2613 consisted mainly of the a-sialon phase coupledwith a small amount of AlN-polytypoid. AlN-polytypoids are AlN defectstructures that result from the incorporation of silicon and oxygen atoms intothe AlN structure (van Tendeloo et al., 1983; Wood and Cheng, 2000). SEMexamination revealed two distinct crystalline phases: the a-sialon phasewith a smooth trait and the AlN-polytypoid phase with speckled features(Fig. 20.1(b)). The speckled appearance of the AlN-polytypoid phase was aresult of the faster etching rate of this phase compared to that of a-sialonwhen NaOH etchant was used (Wood and Cheng, 2000). The a-sialon grainsappeared mainly in an elongated shape with an average diameter of 0.51 mm


and an apparent aspect ratio of 4.1. Image analysis revealed that the areafractions of the AlN-polytypoid phase and the intergranular glassy phase insample CA2613 were ~4% and ~3%, respectively.

20.2.2 Silicon nitride

Two gas pressure sintered (GPS-ed) Si3N4 ceramics, namely SN-F and SN-C, were prepared by the National Industrial Research Institute, Nagoya,Japan. The starting powders were Si3N4 (E-10 grade, ≥95% a-phase, UbeIndustries, Japan), Y2O3 (purity >99.9%, Shinetsu Chemicals, Japan) andAl2O3 (purity >99.9%, Taimei Chemicals, Japan). The composition of thedensification additives was 5 wt% Y2O3 + 2 wt% Al2O3. Material SN-F wasprepared by cold isostatic pressing the powder mixture followed by gaspressure sintering, while material SN-C was fabricated using the same powdermixture with an additional 5 vol% of elongated b-Si3N4 seeds and wasprepared by tape casting, stacking, debinding and gas pressure sintering. Theb-Si3N4 seeds were rod-like single crystal particles with a typical diameterof ~0.5 mm and a length of ~2 mm. Both samples were densified at 1850∞Cfor 6 h under a nitrogen gas pressure of 0.9 MPa. The processing proceduresemployed are described in greater detail in a paper by Hirao et al. (1995).

Microstructures of the two GSP-ed Si3N4 ceramics are shown inFig. 20.2. As can be seen, both materials contain elongated reinforcingb-Si3N4 grains. However, the unseeded sample SN-F (Fig. 20.2(a)) displaysa fine-grained microstructure containing randomly oriented fine elongatedgrains. The distribution of its grain diameter appears to lie in a transitionregion between the distinctly bimodal and the broad monomodal distribution.The average grain diameters of the coarse and fine grains are approximately0.3 and 0.6 mm, respectively. The seeded sample SN-C (Fig. 20.2(b)) exhibits

20.1 SEM micrographs of the two-stage sintered sialon samples:(a) CA1005, (b) CA2613. Specimens were etched in molten NaOH at410∞C for 10 s.

1 mm

(a)

1 mm(b)

AIN-polytypoid

a-sialon


a distinct bimodal distribution of grain diameter in which large elongated b-grains are evenly distributed in a matrix of finer b-grains and an amorphousgrain boundary phase. More significantly, these large elongated b-grains, orwhiskers, appear to lie mainly in the tape casting plane and are oriented inthe casting direction. The average diameters of the large elongated grainsand fine matrix grains are approximately 2 mm and 0.3 mm, respectively.

20.2.3 Silicon carbides

The two SiSiC materials, namely SiC-C and SiC-S (supplied by ConcordEngineering, Australia and Schunk, Germany, respectively) are two-phaseceramics which consist of high-purity SiC and Si. Fig. 20.3 shows details of

20.2 Microstructures of the two GPS-ed silicon nitride materials: (a)SN-F (not seeded, cold pressed), (b) SN-C (seeded, tape cast). Plasmaetching highlights the epitaxial growth of b-sialon on b-Si3N4 cores(indicated by the arrows).

(a) (b)

5 mm 10 mm

Casting direction

(b)

20 mm

(a)

20 mm

20.3 Microstructures of the two siliconized silicon carbide materials,(a) SiC-C and (b) SiC-S, observed with reflected light under an opticalmicroscope, showing different reflective indexes between Si (light)and SiC (dark).


the microstructures of polished sections of the as-received materials imagedin an optical microscope using reflected light. The light Si phase in Fig. 20.3is due to the higher reflectivity of Si in comparison to that of SiC (Lee andRainforth, 1994). As can be seen, both materials possess a duplex microstructurewith angular shaped SiC grains of a bimodal size distribution evenly dispersedin a matrix of fine b-SiC, formed from reaction of the carbon with liquid Si,and free silicon. The average grain size is approximately 50 mm and 6–10mm for large and small SiC grains, respectively, in material SiC-C. Thecorresponding data for material SiC-S are 30 mm and 3–4 mm. The volumefractions of large SiC grains, small SiC grains and free Si, as determinedusing image analysis, are approximately 49%, 35% and 16%, respectively,for material SiC-C, and 58%, 31% and 11%, respectively, for material SiC-S. Note that the volume fraction of small SiC grains includes the originalfine-grained a-SiC particles as well as newly reacted elongated b-SiC grains.

The microstructure of the SiSiC material reflects an interesting processinghistory. It involved the mixing of a-SiC particles, usually with a bimodalsize distribution, with carbon and a thermosetting resin to form a green body.The green compact was then charred to carbonize the resin binder and todrive off the volatiles. Finally, the resulting porous body was infiltrated withmolten Si at temperatures greater than 1500∞C under either vacuum or aninert atmosphere. Liquid Si penetrated the porous body by capillary forceand reacted with the carbon to form fine-grained b-SiC grains, epitaxial b-SiC deposits on the a-SiC grains as well as large b-SiC grains (Lee andRainforth, 1994). The reacted SiC along with the residual Si bonded thebody together to form a final product with good strength.

20.3 Material characterization

20.3.1 Property evaluation

The bulk densities of all the materials were determined using Archimedes’method (AS 1774.5, 1979). The Vickers indentation technique was used tomeasure the hardness in each case. The applied load in the Vickers hardnesstests was 10 kg for silicon nitrides and sialons. However, using the same loadproduced severe lateral cracking in silicon carbides around indents, whichprevented the accurate measurement of the diagonals of indents. Thereforethe load was reduced to 0.3 kg for silicon carbide samples.

Fracture toughness of these materials was also estimated using the Vickersindentation technique by measuring the well-developed radial cracks emanatingfrom the four corners of the indent. The indentation load was 0.3 kg forsilicon carbides, 10 kg for sialons, and 20 kg for silicon nitrides. The reasonfor using different loads for the different materials was to produce well-developed radial cracks of length 2c which were twice as long as the diagonal


length 2a of the indent (Anstis et al., 1981). In addition, since the orientationof the large elongated b-grains, or whiskers, in sample SN-C is highlydirectional, being approximately aligned in the tape casting direction (Fig.20.2(b)), toughness measurements were carried out in directions both paralleland perpendicular to the orientation of the whiskers. It is important to notethat in many of the test materials, fracture toughness is not unique, but ratherdisplays a strong dependence on the crack size, the so-called T-curve or R-curve behaviour (Lawn, 1993). The values reported here thus correspond tothe plateau toughness of the T-curve. The physical and mechanical propertiesof the materials studied in this work are presented in Tables 20.1 and 20.2.

20.3.2 Erosion tests

Samples for erosion tests were cut and machined to dimensions approximately24 mm diameter ¥ 5 mm (sialons), 25 mm ¥ 25 mm ¥ 5 mm (Si3N4), and25 mm ¥ 25 mm ¥ 10 mm (SiC). The surfaces to be eroded were preparedby finish grinding with 800-mesh SiC abrasive paper.

Erosion tests were performed at room temperature in a gas-blast typeerosion test rig described in detail elsewhere (Zhang et al., 2000). Mild steelwas employed as the control material in each test. The test conditions usedare as follows:

Sample to nozzle distance: 13.8 mm;Particle velocity 20 m s–1;Erodent particles SiC (particle size, d50 = 388 mm)Impact angles 30∞ and 90∞Dosage ~200 g for 30∞ impact, ~100 g for 90∞ impact

Table 20.1 Physical and microstructural properties of target materials

Material Density* Porosity§ Grain diameter(kg/m3) (vol%) (mm)

CA1005, Ca a-sialon, PLS-ed + HP-ed 3150 ~2 0.52CA2613, Ca a-sialon, PLS-ed + HP-ed 3208 <1 0.51SN-F, Silicon nitride without seeds 3247 <1 0.3/0.6, nearly

bimodalSN-C, Silicon nitride with seeds 3238 <1 0.3/2, bimodalSiC-C, Reaction bonded silicon carbide 3020 0 8/50, bimodalSiC-S, Reaction bonded silicon carbide 3101 0 4/30, bimodal

*The bulk density of target materials was measured using the water immersionmethod.§The porosity of the sialon ceramics was determined using a combination of thewater immersion method and the image analysis technique. The porosity of therest of the materials was either obtained from the manufacturers data sheet (SiC-Cand SiC-S) or quoted from literature (SN-F and SN-C) (Becher et al., 1998).

Ceram

ic matrix com

posites542

Table 20.2 Mechanical properties of target materials

Sample Hardness Toughness Young’s modulus Four-point flexure Comments(GPa)a (MPa m1/2)b (GPa)c strength (MPa)c

CA1005 18.6 ± 0.4 4.3 ± 0.4 235–245 — Mainly equiaxed grains, aspect ratio ~2CA2613 18.3 ± 0.2 5.6 ± 0.4 235–245 — Elongated grains, aspect ratio ~4SN-F 16.4 ± 0.3 5.9 ± 0.1 304 ~1000 Elongated grainsSN-C 16.7 ± 0.2 7.9 ± 0.5¶ 312 ~1400† 12 MPa m1/2(KIC, SEPB method)¶

5.9 ± 0.2§ ~700‡ 7 MPa m1/2(KIC, SEPB method)§

SiC-C 20.5 ± 0.3 2.4 ± 0.6 350–400 230–300 Mainly equiaxed grainsSiC-S 22.7 ± 0.1 2.4 ± 0.3 300–390 200–300 Mainly equiaxed grains

aThe applied load in Vickers hardness tests was 0.3 kg for silicon carbides, 10 kg for silicon nitrides and sialons.bThe applied load in fracture toughness determination was 0.3 kg for silicon carbides, 10 kg for sialons, and 20 kg for silicon nitrides.cThe values of Young’s modulus and four-point flexure strength of target materials were obtained from literature and suppliers.¶Toughness in the direction perpendicular to the orientation of the large elongated b-grains (whiskers).§Toughness in the direction parallel to the orientation of the large elongated b-grains (whiskers).†When tensile stress was applied parallel to the tape casting direction (whisker orientation).‡When tensile stress was applied perpendicular to the tape casting direction (whisker orientation).


Mass loss from the samples was measured using an analytical balance withan accuracy of ±0.1 mg. Wear volume was calculated from the mass loss andthe bulk density of each material. Cumulative volume loss was plotted as afunction of the amount of erodent impacting on the surface. The steady stateerosion rate, defined as the volume loss from the specimen per unit mass oferodent used, was determined from the slope of the linear part of the plot ofvolume loss against mass of erodent.

The eroded surfaces of all target materials were examined using a JEOLFE6300 scanning electron microscope (SEM) equipped with a field emissiongun. Prior to examination, the specimens were ultrasonically cleaned inethanol for 5 minutes and then sputter coated with carbon to prevent chargeaccumulation on the samples during examination. The accelerating voltageused was 10 kV. To facilitate comparison of damage sustained by differentsamples under different impingement angles, the micrographs, unless specified,were taken at or near the center of the erosion crater.

20.4 Erosion response

20.4.1 Erosion performance

The steady state erosion rates of the target materials followed by SiC erosionat 30∞ and 90∞ impact are shown in Fig. 20.4. Erosion rates for all sampleswere higher for 90∞ impact compared with 30∞. This is consistent with the

20.4 Steady state erosion rates of all target materials eroded by SiCparticles at impingement angles of 30∞ and 90∞. Solid bars labeled Iand II indicate the erosion rates of material SN-C corresponding to30∞ erosion in the direction perpendicular and parallel, respectively,to the whisker orientation.

30∞90∞

CA1005 CA2613 SN-F SN-C SiC-C SiC-S

I II

Ero

sio

n r

ate D E

(m

3 /kg

) ¥

1010

120

100

80

60

40

20

0


expected trend for brittle materials (Sheldon and Finnie, 1966). The rankingof the materials, in descending order of erosion resistance, was the Ca a-sialon ceramics, the two types of self-reinforced silicon nitride materials,and the two SiSiC materials, irrespective of the impact angle. Note that, forthe two self-reinforced Si3N4 materials, SN-C exhibited a higher erosion ratethan SN-F, while for the two sialon ceramics, material CA1005 possessed ahigher erosion rate than CA2613.

In order to elucidate the effect of the whisker orientation on the erosionbehavior of material SN-C, erosion tests were carried out in directions bothparallel and perpendicular to the whisker orientation. It is apparent that inthe highly directional whisker-reinforced silicon nitride material, solid particleerosion in the direction parallel to the whisker orientation resulted in a fasterrate of material removal compared to that in the perpendicular direction (Fig.20.4).

20.4.2 Examination of eroded surfaces

Sialon

The representative features of the eroded surfaces of the two sialon ceramicsare shown in Figs 20.5–20.7. The damaged surfaces of sample CA1005generated by erosion using SiC erodent at normal impact are presented inFig. 20.5. The damage patterns consisted mainly of grain ejection with somelimited plastically deformed materials, indicating that the dominant mechanismof material removal involved intergranular microfracture leading to dislodgmentof individual grains.

The micrographs in Fig. 20.6 show the types of damage produced byerosion at normal impact in sample CA2613. Fig. 20.6(a) depicts the generalview of eroded surfaces of sample CA2613, while Fig. 20.6(b) is the highermagnification views of regions containing faceted grain morphology. It can

1 mm

20.5 SEM micrograph of steady state erosion surface of sampleCA1005 after erosion with SiC at normal impact.


be seen that the damaged surface of sample CA2613 contained mainlyplastically smeared materials with some small-scale grainy regions (Fig.20.6(a)). High-magnification SEM observation revealed that the process ofmaterial removal in sample CA2613 involved the propagation and intersectionof microcracks of grain-facet dimensions incorporating both transgranularand intergranular fracture (indicated by arrows in Fig. 20.6(b)).

The micrograph in Fig. 20.7 shows the types of damage produced byerosion using SiC erodent at 30∞ impact in sample CA2613. As can be seen,damage features consisted mainly of plastically deformed materials originatingfrom the cutting and ploughing actions of the hard, sharp SiC particles.Isolated brittle-fracture regions, characterized by faceted grainy morphology,were also observed.

4 mm

20.7 Surface morphology of sample CA2613 following erosion usingSiC particles at 30∞ impact angle. The particle impact direction isfrom top to bottom of the micrograph.

(b)(a)

1 mm

20.6 Surface morphology of sample CA2613 following erosion usingSiC particles at 90∞ impact angle: (a) low-magnification generalmicrostructure of the eroded surface; (b) higher-magnification imageof regions with grainy morphology.

5 mm


Silicon nitride

Erosion damage in materials SN-F and SN-C has been investigated in detailin a previous paper (Zhang et al., 2005) to which interested readers arereferred. Here we outline the critical findings. It was found that, in materialSN-F, damage features induced by 30∞ and 90∞ impact were similar; bothindicated brittle fracture and plastic deformation of the materials, except thatthe 30∞ impacts produced a higher area fraction of plastically deformedzones with long ploughing trajectories. In the brittle-fracture region,intergranular fracture and transgranular fracture as well as the smearing ofexposed surface grains were observed. In addition, some wear debris, typicallysubmicrometer in size and irregular in shape, were seen on the worn surface,suggesting that they were derived either from the fine, more equiaxed grainsor from the fragmentation of the large, more elongated grains.

Upon repeated impacts, the surface grains along with some loosely bondedwear debris were crushed and smeared to form highly deformed flakes. Atthe same time, crack networks also developed in the target material. Subsequentimpacts resulted in fracture of the underlying silicon nitride grains as well asthe spalling or fragmentation of the smeared flakes, leaving a brittle-fracturezone with faceted morphology.

Damage patterns observed in material SN-C showed a strong dependenceon the impact direction at 30∞ erosion. In the direction parallel to the tapecasting direction (Fig. 20.8(a)), damage sustained by the reinforcing whiskerswas mainly intergranular fracture with a few incidences of transgranular fracture,resulting in pullout of the whiskers, or at least large portions of the whiskers.In addition, microfracture of the fine matrix grains was also seen. However,in the direction perpendicular to the tape casting direction (Fig. 20.8(b)),

1 mmB

A (b)(a)

1 mm

20.8 Worn surfaces of material SN-C after 30∞ SiC erosion withdifferent impact directions. Arrows indicate the tape castingdirection, while the particle impact direction is from top to bottom ofthe micrographs. (a) Parallel to the tape casting direction; (b)perpendicular to the tape casting direction. Arrows labeled A showthe transgranular fracture, while arrows labeled B indicate theintergranular fracture (after Zhang et al., 2005).


more wear debris was observed than that resulting from erosion in the paralleldirection. Although pullout still remained one of the dominant damage featuressustained by the reinforcing whiskers, an increased incidence of microchippingof the whiskers as a result of the impact was clearly evident. Many parallelwhiskers were fractured into a number of sections in the area of contact withthe impinging particles, exhibiting a severely crushed morphology consistingof many fragments.

The eroded surface resulting from 90∞ impact exhibited a higher areafraction of brittle-fracture zones compared to that produced by 30∞ impacts.Pullout and transgranular chipping of the whiskers as well as microfractureof the fine matrix grains were observed.

Silicon carbides

In gas-blast erosion, solid particles are accelerated along a parallel wallednozzle, but exit the nozzle as a diverging plume. The distribution of particletrajectories within the plume is well defined (Shipway, 1997), and theprobability of a particle striking the target becomes smaller as the angle ofparticle trajectory with respect to nozzle axis increases. Thus the examinationof the edge of the erosion crater permits the investigation of the damageinduced by single particle impact.

The early stage erosion surface of material SiC-C resulting from SiCerosion at 90∞ impact is shown in Fig. 20.9. A few isolated pits similar tothose shown in Fig. 20.9(a), marked by arrows, were found in these regions.Interestingly, detailed examination revealed that damage features associatedwith individual impact sites were quite different. In the impact crater marked

(b)(a)

10 mmA

B

10 mm

Si

SiC Si

20.9 SEM micrographs showing surface morphology of isolatedparticle impacts on material SiC-C following 90∞ SiC erosion: (a)secondary electron (SE) image; (b) back-scattered electron (BSE)image where atomic number contrast between Si (light) and SiC isperceived. The images were taken near the edge of the erosioncrater.


by the arrow labeled A, classical chipping or conchoidal fracture resultingfrom lateral cracking was evident. In contrast, in the nearby crater labeled B,dislodgment of fine SiC grains was observed.

A back-scattered electron (BSE) image of the same area in Fig. 20.9(a) isshown in Fig. 20.9(b) where the bright phase indicates the regions rich in Si.The light grey phase presents the two-phase SiC–Si region while the darkgrey phase, highlighted by the frame, shows the large SiC grains. As can beseen from Fig. 20.9(b), conchoidal fracture from lateral crack chipping occurredwithin individual coarse SiC grains, while the grain dislodgment took placein the two-phase region. In addition, small-scale chipping damage was alsoobserved on the surface of large SiC grains, suggesting that not every impactingparticle could effectively produce a significant lateral crack. This is probablydue to the threshold effect on lateral crack initiation where only those particleswith high kinetic energy are capable of producing lateral cracking or evenmultiple impacts are necessary to accumulate the requisite levels of stress togenerate lateral cracking.

The steady state erosion surface of material SiC-C produced by SiC erosionat 90∞ impact is shown in Fig. 20.10. The worn surface was much rougherthan that resulting from early stage erosion. Areas of smeared-looking ordeformed material were also observed in addition to the brittle-fracture regions(Fig. 20.10(a)). Furthermore, in the brittle-fracture zones, two distinct typesof damage were present: (1) large-scale transgranular chipping resultingfrom the link-up of lateral cracks; and (2) dislodgment of the fine SiC grainsand occasional small-scale transgranular chipping of the elongated b-SiCgrains, highlighted by elliptical frames in Fig. 20.10.

The BSE image of the same area in Fig. 20.10(a) is shown in Fig. 20.10(b).The effect of microstructure on erosion damage is immediately evident. Theerosion damage of large SiC grains was predominately transgranular chipping

(a) (b)

10 mm 10 mm

SiC

20.10 SEM micrographs showing surface morphology of materialSiC-C following 90∞ SiC erosion: (a) SE image, (b) BSE image. Theimages were taken near the center of the erosion crater, wheresteady state erosion occurred.


of lateral cracks (see the dark grey regions), indicated by the frame withdashed lines, in Fig. 20.10(b). On the other hand, the damage features observedin the two-phase region were grain dislodgment, limited small-scaletransgranular chipping as well as highly plastically deformed materials, asseen in regions containing the scattered bright Si phase in Fig. 20.10(b). Thedeformed material is believed to contain free Si as well as fine SiC debrisand thus should have a bright appearance relative to SiC grains under theBSE mode. However, such differences are often diminished due to the rougherosion surface and the non-uniform distribution in thickness of the surfacedeformed materials.

Steady state erosion surfaces of material SiC-S generated by SiC particlesat 30∞ and 90∞ impact are shown in Figs 20.11(a) and 20.11(b), respectively.Remarkably, damage features were very similar in the two cases, both involvingtransgranular chipping associated with lateral cracks and intergranular fractureof fine SiC grains as well as plastic smearing of the deformed materials.Additionally, there was a noticeable lack of ploughing marks on the 30∞eroded surface, suggesting that the plastic cutting mechanism was suppressedeven when the hard, sharp SiC erodent was used.

The Vickers indentation impressions of the two SiSiC materials, at least at10 kg load or greater, were surrounded by regions of widespread lateralcracking and associated surface chipping (Fig. 20.12(a)). Figure 20.12(b) isa BSE image showing the damaged surface located at the edge of an indentationimpression in material SiC-S, while Fig. 20.12(c) is a higher-magnificationview of Fig. 20.12(b) illustrating a typical trajectory of the indentation inducedcrack. As can be seen, the advancing crack tip propagates through the SiCgrains with no sign of deflection or bridging, indicating a low fracture toughnessof these materials (Table 20.2).

(b)(a)

10 mm 10 mm

20.11 Steady state erosion surfaces of material SiC-S following SiCerosion at (a) 30∞ impact, and (b) 90∞ impact. The particle impactdirection for 30∞ impact is from top to bottom of the micrograph.


20.5 Microstructure and mechanical properties

It has been long recognized for Si3N4- and SiC-based materials that thereexists a close relationship between microstructural parameters and mechanicalproperties (Becher et al., 1993; Padture and Lawn, 1994; Hoffmann, 1994;Zhao et al., 1997; Chen and Rosenflanz, 1997; Becher et al., 1998; Ellen etal., 1998; Kleebe et al., 1999). Compared to Si3N4 and SiC, sialon ceramicshave greater phase complexity and more degrees of freedom for tailoring ofmicrostructures and consequently mechanical properties. Thus, the possiblemicrostructural tailoring of sialon ceramics means that these materials offeran excellent modeling system in any investigation of the effect of variousmicrostructural aspects on the mechanical properties and the erosion behaviorof ceramic materials.

20.12 SEM micrographs of Vickers indentation-induced damage inmaterial SiC-S (peak load 10 kg): (a) SE image, (b) and (c) BSEimages.

50 mm(a)

(b)30 mm

15 mm(c)


The microstructural–mechanical properties relation of the Ca a-sialonmaterials has been studied in detail previously (Zhang et al., 2001; Zhangand Cheng, 2003). Some interesting trends were observed. It was found thatthe hardness of the materials decreases as the x-value increases. The decreasein the amount of a-sialon phase and the increase in the intergranular glasscontent with increasing x-value may account for this. It was also found thata high a-sialon content coupled with low intergranular glass and pore contentsgave an optimized hardness (Zhang et al., 2001; Zhang and Cheng, 2003). Incontrast, the apparent aspect ratio of the a-sialon grains appeared to havelittle influence on the hardness value of these materials. As tabulated inTable 20.2, despite a dramatic difference in the grain aspect ratio betweenthe HP-ed samples CA1005 and CA2613, the hardness values of the twomaterials are virtually the same. It may be argued that the relatively highporosity in sample CA1005 can result in a decline in hardness; however, thehigher glass content in sample CA2613 is also known to inhibit hardness.Considering that the porosity and the glass content in these materials arequite low, being less than 2 and 3 vol%, respectively, it is reasonable toconclude that the grain aspect ratio in a-sialon ceramics has little impact ontheir hardness values. The present finding is consistent with previousobservations where a-sialons containing various amounts of elongated grainswith very different aspect ratios possessed almost identical hardness values(Chen and Rosenflanz, 1997; Kim et al., 2000).

On the other hand, a coarse grain size and a high aspect ratio can both giverise to the fracture toughness. The effect of grain aspect ratio on the fracturetoughness is evidenced in HP-ed samples CA1005 and CA2613 where fracturetoughness increases as the grain aspect ratio increases. The toughening effectobserved in this study is mainly attributed to the crack bridging mechanism.However, to further improve the toughness, cracks must propagate alonggrain interfaces rather than through the grains. This interface debondingprocess appears to be governed by the chemistry of the oxynitride glass atthe grain boundaries (Hoffmann, 1994; Kleebe et al., 1999; Becher et al.,1994).

The chemistry of the grain boundary glass probably holds the key tounderstanding why the fracture toughness of a-sialon ceramics is significantlylower than that of silicon nitride. Current best room-temperature values offracture toughness for a-sialon and Si3N4 ceramics were ~6 MPa m1/2 (Zhaoet al., 1997; Chen and Rosenflanz, 1997) and >10 MPa m1/2 (Table 20.2),respectively. The Nd- and Y-stabilized a-sialon ceramics fabricated by Chenand Rosenflanz (1997) and later Kim et al. (2000) exhibited a microstructureconsisting of large elongated a-sialon grains imbedded in a matrix of fine-grained a-sialon. This microstructure is very similar to that of Si3N4 whichwas described to have the best properties where large elongated b-sialongrains were evenly dispersed in a matrix of fine-grained b-Si3N4 and an


amorphous grain boundary phase. However, the fracture toughness of thesea-sialons attained only half of the fracture toughness value of silicon nitrideceramics. One possible reason is that in a-sialon, unlike in b-sialon and b-Si3N4, both grain boundary glass and a-sialon grains contain identical elements,resulting in a strong bonding between a-sialon grains and the glassy matrix.Therefore to further improve the fracture toughness of a-sialon ceramics,investigations into compositional design of the intergranular glass that leadsto a weakened bonding strength between a-sialon grains and the glassymatrix is necessary.

20.6 Microstructure and erosion mechanisms

Traditionally, hardness and toughness are often used to model the erosionbehavior of ceramic materials (Evans et al., 1978; Wiederhorn and Lawn,1979). With the development of indentation fracture mechanics, two elastic–plastic theories have been developed to model the erosion behavior of brittlematerials. Evans et al. (1978) considered the dynamic elastic–plastic responseof an impacting particle, while Wiederhorn and Lawn (1979) assumed theparticle-target contact velocity was slow (relative to sonic velocity), i.e. thequasi-static condition, and the kinetic energy of the erodent particle wasabsorbed completely by the plastic flow of the target. However, both theoriesassumed that the lateral cracks were responsible for material removal andboth predicted a power-law dependence for the erosion rate, DE, on erodentand target properties given by

DE µ vnD2/3r p K H qC–4/3 (20.2)

where v, D and r are the velocity, mean size and density of the erodentparticles, and KC and H are the fracture toughness and hardness of the targetmaterial. The exponents n, p and q differ in the two models, being 3.2, 1.3and –0.25 for the dynamic and 2.4, 1.2 and 0.11 for the quasi-static model,respectively. The models indicate that the erosion rate of ceramic materialsshould have a strong inverse dependence on the fracture toughness, but amuch weaker dependence on the hardness of the material. However, such aprediction seems to be contrary to the present findings. Mechanical propertyevaluations showed that material SN-C exhibited a significantly highertoughness, in the direction perpendicular to the tape casting direction, thanmaterial SN-F, while the hardness of the two materials was almost identical(Table 20.2). According to the theoretical models (eq. (20.2)), material SN-C should have a better erosion resistance, especially when eroded in a directionperpendicular to the tape casting direction, than material SN-F. Nevertheless,erosion tests showed that material SN-F exhibited better erosion resistancethan material SN-C under both 30∞ and 90∞ impacts (Fig. 20.4).


20.6.1 Sialon

Erosion tests showed that sample CA1005 exhibited a higher material removalrate than sample CA2613 at both 30∞ and 90∞ impacts. SEM examination ofsurfaces of the two samples generated by erosion at 90∞ impact revealed thatin material CA1005 the dominant material removal mechanism was grainejection caused by grain boundary microcracking, while in material CA2613the mechanisms involve combined intergranular fracture of low-aspect-ratiograins and transgranular chipping of large interlocking grains.

Consider these results in terms of the different microstructures in the twomaterials. Samples CA1005 and CA2613 were fabricated from differentbatches and were both PLS-ed first and then HP-ed at the high temperatures.A low content of intergranular glassy phase, or liquid phase above the eutectictemperature, can hinder the densification process of composition CA1005.As a result, sample CA1005 possesses higher porosity, by a factor of ~3,than sample CA2613 – a composition containing a considerably higher amountof intergranular glass. In material CA1005 a low intergranular glass contentcoupled with a high porosity can result in a weakened bonding betweengrains. In addition, pores, especially sintering pores at grain boundary triplepoints, can be regarded as a stress-concentrating flaw system and hence arepreferred sites for microcrack initiation (Lawn, 1993). Therefore, upon impact,microcracks readily initiate at the multi-grain junctions and consequentlypropagate along the two-grain interfaces and link up to from crack networks.The fine, equiaxed structure of material CA1005 offers limited resistance tocrack propagation. As a result, material is removed from the target surfacevia severe grain dislodgment (Fig. 20.5).

In material CA2613, an elongated grain morphology coupled with a smallamount of grain boundary glass can give rise to enhanced erosion resistance.Grain boundary glass fills into the pores and facilitates the particlerearrangement during the sintering and also has the potential to absorb orcushion the stress induced from solid particle impact via viscous flow. Theelongated grain morphology can have at least two beneficial effects on erosionresistance of ceramic materials (Zhang et al., 2001): it can enhance crackdeflection and bridging by blocking the crack path, forcing the crack todetour; and it can hinder grain dislodgment due to an interlocking effect.Both effects suggest that more energy is required to propagate and link upcracks between neighboring impact sites to form crack networks and toremove the interlocked grains from the target surface.

20.6.2 Silicon nitride

Detailed discussion in respect to the erosion behavior of self-reinforced orin-situ reinforced silicon nitride materials can be found elsewhere (Zhang etal. 2005). Here we summarize the important arguments.


In the erosion of silicon nitride with large directional reinforcing whiskers,the reinforcing whiskers are seen to undergo pullout and fragmentation.More significantly, the material removal mechanism of the whiskers is foundto vary with the solid particle impacting direction, i.e. in a direction parallelor perpendicular to the tape casting direction. When solid particles attack thesurface in the direction parallel to the casting direction, the whiskers areremoved primarily by debonding, pullout, or fracture of a large portion ofthe whisker. Conversely, when erodent particles strike the surface in thedirection perpendicular to the casting direction, the whiskers experience anincreased incidence of multiple fractures in addition to pullout. In manycases, the multiple fractured sections are chipped away or crushed by theimpacting particles. Since erosion in the direction perpendicular to the whiskerorientation entails multiple fracture and chipping, which requires higherenergy compared to the whisker-dislodgment in the parallel direction, theerosion loss is less in the direction perpendicular to the whisker orientationthan that parallel to the whisker orientation.

The present study indicates that self-reinforced silicon nitride ceramicsdisplaying a high steady-state toughness might not exhibit good erosionresistance. This finding may be explained by the following reasons. In afracture toughness test, when the propagating crack tip interacts with thereinforcing whiskers, the weak interface between the whiskers and the matrixpromotes crack bridging, resulting in a crack-size dependent toughness. Thisis particularly the case for SN-C material where highly directional, largeelongated grains are well dispersed in a fine, submicrometer grain sizedmatrix. Such a microstructure ensures a high steady-state fracture toughnessowing to the large number of reinforcements and the large length of thedebonded interfaces between the whiskers and the matrix.

In an erosion test, the situation where a crack grows at equilibrium overa long distance and encounters a number of bridging grains does not exist.Rather, multiple microcracks develop simultaneously in the vicinity of theimpact site and linking up of microcracks between the neighborhood impactsites forms a network of short cracks, leading to material loss from the targetsurface. The weak interface, which is essential for the R-curve behavior, canfacilitate the formation of interfacial microcracking and thus result in thedislodgment of the large reinforcing whiskers, leading to a high rate ofmaterial removal. More significantly, reinforcing whiskers for SN-C are welldispersed and oriented parallel to the eroding surface. During erosion, thewhiskers tend to get plucked out or sheared out of the matrix rather thanundergo the partial pullout which leads to crack bridging. On the other hand,the randomly oriented reinforcing grains in material SN-F provide aninterlocking effect. Upon erosion, the interlocking hinders the dislodgmentof the large reinforcing grains, resulting in a relatively low material removalrate.


It has been reported in the literature (Srinivasan and Scattergood, 1991;Marrero et al., 1993) that when fracture toughness values corresponding toshort cracks relevant to the erosion phenomenon are used, better correlationbetween the fracture toughness and the erosion results can be obtained.While the R-curve behavior or the operative short-crack toughness may givesome indications of the erosion response of brittle materials, the use of short-crack toughness to quantitatively predict the erosion rate must be made withcaution. The key features for a pronounced R-curve behavior are largedirectional reinforcing whiskers coupled with weak interfaces (Becher et al.,1998). The present findings clearly point out that both directional whiskersand a weak interface could deteriorate the erosion resistance owing to thesubstantial dislodgment of the large reinforcing grains.

20.6.3 Silicon carbides

The erosion mechanism in SiSiC materials apparently involves the linkingup of lateral cracks, dislodgment of the fine SiC grains, occasional small-scale transgranular fracture of the fine, elongated b-SiC grains, and plasticsmearing of the free Si phase and fine debris. The dominant mechanisms ofmaterial removal are, however, conchoidal fracture within individual largeSiC grains and dislodgment of the fine-grained SiC particles in the two-phase region.

Despite the considerable differences in microstructure of the two SiSiCmaterials, e.g. the size and volume fraction of the bimodal SiC grains as wellas the amount of the free Si phase, the erosion rates of the two materials varyby less than 10% for both 30∞ and 90∞ impacts. The lack of a microstructuraleffect on erosion rate under the current experimental conditions suggeststhat although there exists a distinct difference in mechanisms of materialremoval between the large SiC grains and the two-phase regions, the actualrate of material removal may be very similar. Indeed, SEM observations ofthe eroded surfaces showed a uniform type of surface morphology; thereis no noticeable preferential wear of particular regions (Figs 20.10 and20.11).

A previous study reported in the literature showed that pure Si erodesmuch faster, by approximately two orders of magnitude, than single-phaseSiC (Routbort et al., 1980). However, preferential erosion of the two-phaseSiC–Si region was not observed in the current investigation. This can beattributed in part to the fact that the fine SiC particles and the newly reactedelongated b-SiC grains act as reinforcement to the residual silicon phase,which has effectively improved the erosion resistance of the two-phase regions,and in part to the severity of the SiC erosion which can effectively causelateral chipping of the large SiC grains. In fact, previous studies of erosionof SiSiC materials using fine Al2O3 erodent particles (Routbort et al., 1980)


and coal slurry (Shetty et al., 1982) have observed the preferential erosion ofthe two-phase region resulting in a surface morphology exhibiting mesa-likeprotrusions of the large SiC grains. This is because both Al2O3 and coal aremuch softer than SiC and therefore are not capable of causing significantfragmentation of the large SiC grains. As a result, the two-phase SiC–Siregion between the large SiC grains erodes relatively fast, leaving the flat-topped mesa-like protrusions of the large grains separated by valleys of thetwo-phase region (Routbort et al., 1980).

Finally, no clear evidence of dislodgment of the large SiC grains wasrevealed in the two SiSiC materials under the current experimental conditions,suggesting that strong bonding exists between the large SiC grains and themixed SiC–Si phase. This scenario is further supported by SEM observationsof the interaction between a propagating crack and SiC grains of these materials,where the advancing crack tip cuts through the large SiC grains with no signof deflection or bridging (Fig. 20.12(c)).

20.7 Conclusions

The mechanical properties of Ca a-sialon ceramics were found to be closelyassociated with their microstructures. It was observed that hardness wasdependent on a combination of a-sialon content, grain size, porosity and theamount of glassy phase in the material, while the fracture toughness increasedwith the increasing grain size and grain aspect ratio. However, to furtherimprove the toughness of these materials, a weakened interface between thea-sialon grains and intergranular glass is necessary.

The Ca a-sialon and silicon nitride ceramics in the current study bestillustrate the effect of microstructural features, such as grain size, grainmorphology and grain boundary glass, on the erosion response of ceramicmaterials. The present results showed that materials consisting of randomlyoriented, elongated grains coupled with a small amount of intergranularglass, i.e. CA2613 and SN-F, exhibited the best erosion resistance. This ismainly due to the fact that the elongated grain morphology gives rise to abetter resistance to crack propagation/networking and has an interlockingeffect that prevents material removal via grain dislodgment. On the otherhand, an optimum amount of grain boundary glass can also improve theerosion resistance of ceramic materials by absorbing the impact stress viaviscous flow as well as by enhancing the bonding strength between thegrains and the matrix.

The highly directional whisker-reinforced silicon nitride material, SN-C,showed a lower erosion resistance than the finer grained self-reinforcedsilicon nitride material, SN-F, and the Ca a-sialon material, CA2613. Thiscan be attributed to substantial pullout of the reinforcing whiskers during theerosion of material SN-C. The significant pullout of the whiskers is a direct


consequence of the highly directional orientation of the whiskers, which ispractically parallel to the erosion surface, and a weak interface between thewhiskers and the matrix.

The equiaxed sialon material CA1005 displayed a predominantlyintergranular mode of fracture, associated with a relatively high erosion ratecompared to its elongated counterpart CA2613. This is because compositionCA1005 contains a very small amount of intergranular glass, leading tohigher porosity and consequently a weak grain boundary in these materials.The weak grain boundary promotes grain boundary cracking under the externalstress. In addition, the almost equiaxed grain morphology of these materialshas little resistance to grain dislodgment. As a result, the dominant erosionmechanism of composition CA1005 is grain ejection.

The relatively low erosion resistance of the two SiSiC materials incomparison to silicon nitride and sialon ceramics is due in part to the lateralchipping of the large SiC grains and in part to the relatively low erosionresistance of the soft Si phase.

20.8 References

Anstis, G.R., et al. (1981), ‘A critical evaluation of indentation techniques for measuringfracture toughness: I, Direct crack measurements’, J. Am. Ceram. Soc., 64(9), 533–8.

AS 1774.5 (1979), ‘The determination of density, porosity and water absorption’, StandardsAssociation of Australia, 1–3.

Becher, P.F. et al. (1993), ‘The influence of microstructure on the mechanical behaviourof silicon nitride ceramics’, in Chen I-W, Silicon Nitride – Scientific and TechnologicalAdvances, MRS Symposium Proceedings, 287, MRS, Pittsburgh, PA, p. 147.

Becher, P.F. et al. (1994), ‘Microstructural contribution to the fracture resistance ofsilicon nitride ceramics’, in Hoffmann M. J. and Petzow G., Tailoring of MechanicalProperties of Si3N4 Ceramics, NATO ASI Series, Series E: Applied Science, 276,Dordrecht, Kluwer Academic, 87–100.

Becher, P.F. et al. (1998), ‘Microstructural design of silicon nitride with improved fracturetoughness: I, Effects of grain shape and size’, J. Am. Ceram. Soc., 81(11), 2821–30.

Bennison, S.J. and Lawn, B.R. (1989), ‘Role of interfacial grain-bridging sliding frictionin crack-resistance and strength properties of non-transforming ceramics’, Acta Metall.Mater., 37(10), 2659–71.

Cao, G.Z. and Metselaar, R. (1991), ‘a-Sialon ceramics: a review’, Chem. Mater., 3, 242–52.

Chen, I.-W. and Engineer, M. (1999), ‘Model for fatigue crack growth in grain-bridgingceramics’, J. Am. Ceram. Soc., 82(12), 3549–60.

Chen, I.-W. and Rosenflanz, A. (1997), ‘A tough SiAlON ceramic based on a-Si3N4 witha whisker-like microstructure’, Nature, 389, 701-4.

Ellen, Y.S. et al. (1998), ‘Microstructural design of silicon nitride with improved fracturetoughness: II, Effects of yttria and alumina additives’, J. Am. Ceram. Soc., 81(11),2831–40.

Evans, A.G. et al. (1978), ‘Impact damage in brittle materials in the elastic–plastic responseregime’, Proc. R. Soc. London, Ser. A, 361, 343–65.


Fairbanks, C.J. et al. (1987), ‘Crack-interface grain bridging as a fracture resistancemechanism in ceramics’, J. Am. Ceram. Soc., 70, 279–94.

Hampshire, S. et al. (1978), ‘a-Sialon ceramics’, Nature, 274, 880–2.Hirao, K. et al. (1995), ‘Processing strategy for producing highly anisotropic silicon

nitride’, J. Am. Ceram. Soc., 78(6), 1687–90.Hoffmann, M.J. (1994), ‘Analysis of microstructural development and mechanical properties

of Si3N4 ceramics’, in Hoffmann, M. J. and Petzow, G., Tailoring of MechanicalProperties of Si3N4 ceramics, NATO ASI Series, Series E: Applied Science, 276,Dordrecht, Kluwer Academic, 59–71.

Jack, K.H. (1976), ‘Review - sialons and related nitrogen ceramics’, J. Mater. Sci., 11,1135–58.

Kelly, J.F. et al. (1991), ‘In situ measurements of bridged crack interfaces in the scanningelectron microscope’, J. Am. Ceram. Soc., 74, 3154–7.

Kim, J. et al. (2000), ‘Microstructure control of in-situ-toughened a-sialon ceramics’, J.Am. Ceram. Soc. 83(7), 1819–21.

Kleebe, H.-J. et al. (1999), ‘Microstructure and fracture toughness of Si3N4 ceramics:combined roles of grain morphology and secondary phase chemistry’, J. Am. Ceram.Soc., 82(7), 1857–67.

Lange, F.F. (1973), ‘Relation between strength, fracture energy, and microstructure ofhot-pressed silicon nitride’, J. Am. Ceram. Soc., 56(10), 518–22.

Lawn, B.R. (1993), Fracture of Brittle Solids, Cambridge, Cambridge University Press.Lee, W.E. and Rainforth, W.M. (1994), Ceramic Microstructures, Property Control by

Processing, London, Chapman & Hall.Marrero, M. et al. (1993), ‘Solid-particle erosion of in situ reinforced Si3N4’, Wear, 162–

164, 280–4.Ohji, T. et al. (1995), ‘Fracture resistance behaviour of highly anisotropic silicon nitride’,

J. Am. Ceram. Soc., 78(11), 3125–28.Padture, N.P. and Lawn, B.R. (1994), ‘Toughness properties of a silicon carbide with an

in situ induced heterogeneous grain structure’, J. Am. Ceram. Soc., 77(10), 2518–22.Routbort, J.L. et al. (1980), ‘The erosion of reaction-bonded SiC’, Wear, 59, 363–75.Sheldon, G.L. and Finnie, I. (1966), ‘On the ductile behaviour of nominally brittle materials

during erosive cutting’, Trans. ASME, 88B, 387–92.Shetty, D.K. et al. (1982), ‘Coal slurry erosion of reaction-bonded SiC’, Wear, 79, 275–

9.Shipway, P.H. (1997), ‘The effect of plume divergence on the spatial distribution and

magnitude of wear in gas-blast erosion’, Wear, 205, 169–77.Srinivasan, S., and Scattergood, R.O. (1991), ‘R curve effects in solid particle erosion of

ceramics’, Wear, 142, 115–33.Swain, M. (1994), Structure and Properties of Ceramics, Weinheim, VCH.van Tendeloo, G. et al. (1983), ‘Characterization of AlN ceramics containing long-period

polytypes’, J. Mater. Sci., 18, 525–32.Wiederhorn, S.M. and Lawn, B.R. (1979), ‘Strength degradation of glass impacted with

sharp particles: I, Annealed surfaces’, J. Am. Ceram. Soc., 62(1–2), 66–70.Wood, C.A. and Cheng, Y.-B. (2000), ‘Phase relationships and microstructures of Ca and

Al-rich a-sialon ceramics’, J. Eur. Ceram. Soc., 20, 357–66.Zhang, Y. et al. (2000), ‘Erosion of alumina ceramics by air- and water-suspended garnet

particles’, Wear, 240, 40–51.Zhang, Y. et al. (2001), ‘Influence of microstructure on the erosive wear behaviour of Ca

a-sialon materials’, J. Eur. Ceram. Soc., 21(13), 2435–45.


Zhang, Y. and Cheng, Y.-B. (2003), ‘Microstructural design of Ca a-sialon ceramics:effects of starting compositions and processing conditions’, J. Euro. Ceram. Soc., 23(9), 1531–41.

Zhang, Y. et al. (2005), ‘Erosion response of highly anisotropic silicon nitride’, J. Am.Ceram. Soc., 88(1), 114–20.

Zhao, H. et al. (1997), ‘Elongated a-sialon grains in pressureless sintered sialon ceramics’,J. Eur. Ceram. Soc., 18, 1053–7.

560

21.1 Introduction

Oxynitride glasses are special types of silicate or alumino-silicate glasses inwhich oxygen atoms in the glass network are partially replaced by nitrogenatoms (Hampshire, 2003a). They occur in M-Si-O-N, M1-M2-Si-O-N, M-Si-Al-O-N and M1-M2-Si-Al-O-N systems where M, M1 and M2 are modifyingcations such as the alkali metals (Li, Na, K), the alkaline earths (Mg, Ca, Ba,Sr) or Y, La and the rare earth lanthanides. Oxynitride glasses exhibit higherrefractoriness, elastic modulus, viscosity and hardness compared with thecorresponding oxide glasses as a result of the extra cross-linking provided bynitrogen within the glass structure.

Oxynitride glasses may be heat treated to form glass-ceramics, effectivelymulti-phase composites. The process involves heat treatment at two differenttemperatures, firstly to induce nucleation, then to allow crystal growth of thenuclei. The crystalline phases formed depend on both the composition of theparent glass and the temperatures used for heat treatment. The extent of theirformation and growth, the relative amounts and distributions of differentphases (including residual glass) and their characteristics will determine theoverall properties of the particular composite. The formation of these typesof materials and their properties is outlined below.

Other investigations have incorporated SiC nano-phase inclusions intooxynitride glass matrices in order to form composites with improved mechanicalproperties.

21.2 Potential applications

Speciality or ‘new’ glasses and glass-ceramics with novel functions are beingused in modern industrial sectors such as optoelectronics, microelectronics,communications technologies, biomedical devices and niche areas of theautomotive and architectural sectors. Oxynitride glasses and glass-ceramicsand their composites expand the range of ‘new’ glasses and, in view of their

21Oxynitride glasses – glass ceramics

composites

S H A M P S H I R E, University of Limerick, Ireland

Oxynitride glasses – glass ceramics composites 561

durability, higher refractoriness, tailored thermal properties, etc., there arethe following areas of potential application:

∑ Passive coatings on electronic substrates∑ Novel glaze systems for refractory protection∑ High-temperature joining of ceramics and of ceramics to metals∑ Higher temperature range glass-ceramics and composites for structural

applications.

21.3 Oxynitride glass/glass-ceramic composites

21.3.1 Oxynitride glasses

Various reviews on oxynitride glasses by Leng-Ward and Lewis (1990),Thompson (1992), Sakka (1995) and Hampshire (2003a) have given insightsinto the understanding of structure–property relationships in these materials.Interest in oxynitride glasses developed following the discovery that siliconnitride-based ceramics contained oxynitride glassy phases at their grainboundaries as a result of the use of oxide sintering additives, which reactwith silica on the surface of the nitride particles and the nitride itself to formliquid phases that cool as intergranular glasses (Hampshire, 1994). Thecomposition and volume fraction of these glass phases determine the propertiesof the material, particularly their high-temperature mechanical behaviour(Hampshire and Pomeroy, 2004).

The extent of the glass-forming regions in various M-SiAlON systems (M= Mg, Y, Ca, Nd, etc.) has been studied (Hampshire et al., 1985) and representedusing the Jänecke prism, with compositions expressed in equivalents insteadof atoms or gram-atoms. Thus, equivalent percentage (e/o) of cation M isgiven by:

e/o M = v [M] 100

v [M] + v [Si] + v [Al]M

M Si Al

¥(21.1)

where [M], [Si] and [Al] are, respectively, the atomic concentrations of M,Si and Al, VM, VSi, VAl are respectively, the normal valencies for M, Si andAl. Similarly, for the anions, the equivalent percentage (e/o) of nitrogen isgiven by:

e/o N = (3[N]) 1002[O] + 3[N]

¥(21.2)

where [O] and [N] are, respectively, the atomic concentrations of oxygenand nitrogen.

Figure 21.1 shows the glass-forming region in the Y-Si-Al-O-N Jäneckeprism (Drew et al., 1981). The region is seen to expand initially as nitrogenis introduced and then diminishes above approximately 10 e/o N until thesolubility limit for nitrogen is exceeded.


Further investigations on oxynitride glass formation, structure and propertieshave shown that oxynitride glasses have higher glass transition temperaturesTg, elastic moduli, viscosities and hardness values compared with thecorresponding oxide glasses (same cation ratio) as a result of the extra cross-linking provided by nitrogen within the glass structure (Hampshire, 2003a).Properties typically increase linearly with nitrogen content. Viscosity increasesby more than 2 orders of magnitude from a value of 1010.3 Pa.s for the oxideglass to 1012.6 (~Tg) as 18 e/o oxygen is replaced by nitrogen at 950∞C.Changes in the cation ratios result in smaller changes in properties thanchanges in N:O ratio.

Ohashi and Hampshire (1991) Ohashi et al., (1995) studied glasses invarious LnSiON systems while Ramesh et al. (1997) investigated a range ofLnSiAlON glasses. For glasses with the same Ln-Si-Al-O-N ratios but differentrare earth lanthanide cations (Ln), properties such as density, hardness, Tg,elastic modulus and viscosity increase linearly with increasing cation fieldstrength (CFS) or decreasing cationic radius. Substitution of one Ln cationby another appears to cause no change in the overall glass structure.

21.3.2 Crystallisation of oxynitride glasses to formmulti-phase glass-ceramic composites

As with other silicate glasses, oxynitride glasses may be heat treated atappropriate temperatures to crystallise as glass-ceramics (Leng-Ward andLewis, 1990; Thompson, 1992; Hampshire, 1993). The crystalline phasesformed depend on both the composition of the parent glass and the heat-

4YN

Si3N4

2Y2O3

4AlN

eq% N

3/2 Si2N2O

*Y2O3–SiO2 eutectic

a

b

6/7 Y2Si2O7

3SiO22Al2O3

21.1 Glass-forming region in the Y-Si-Al-O-N Jänecke prism oncooling from 1700∞C (after Drew et al., 1981).

10

50


treatment process, and the extent of their formation and growth will determinethe properties of the particular material.

The conventional process to produce a glass-ceramic involves two steps:a first-stage heat treatment of the glass at a temperature just above Tg toinduce nucleation, followed by heating to a second higher temperature (thecrystallisation temperature, Tc) to allow crystal growth of the nuclei. Manyglasses require the addition of a nucleating agent to promote the crystallisationprocess (Tredway and Risbud, 1984), but oxynitride glasses appear to beself-nucleating because nucleation occurs heterogeneously on FeSi particlesas found by Dinger et al. (1988) and Besson et al. (1993) in Y2O3-SiO2-AlNglasses.

21.3.3 Glass-ceramic composites in theY-Si-Al-O-N system

Many studies of crystallisation of oxynitride glasses have concentrated onthe Y-Si-Al-O-N system. Thompson (1989) outlined the various crystallinephases that exist in this system as shown in the Jänecke prism in Fig. 21.2.

Leng-Ward and Lewis (1985) studied the crystallisation at 1250∞C of aseries of glasses of constant cation composition (in e/o) 26Y:42Si:32Al with0 to 30 e/o nitrogen. In the oxide glass, yttrium disilicate is formed as theprimary phase on heat treatment. As the nitrogen content increased, gradual

Al4O6m 2050

Al4N4s ~ 2200

YAGm 1930

b¢

a ¢

d > 2000

Y4O6m 2410

Y4N4m > 2700

Y2SiO5m 1980

Y2Si2O7m 1775B

M

Si3O6m 1720

Y2SiAlO5ND ~ 1200

Si2N2OD ~ 1900

Si3N4d ~ 1900

N-melilitem ~ 1900

N-a-wollastonited ~ 1400

N-YAMm 1900-2000

N-apatitem 1700-1800

21.2 Y-Si-Al-O-N Jänecke prism showing crystalline phases and solidsolution ranges.


replacement of yttrium disilicate phase by yttrium aluminium garnet (YAG)was observed and nitrogen was mainly incorporated into Si2N2O.

The glass-ceramic transformations in a glass of composition (in e/o)28Y:56Si:16Al:83O:17N were studied by Ramesh et al. (1998) using bothclassical and differential thermal analysis techniques. These two methodswere found to be in close agreement. Optimum nucleation and crystallisationtemperatures were determined in relation to the glass transition temperature.The major crystalline phases present are mixtures of silicon oxynitride anddifferent forms of yttrium disilicate which exists as a, b, g , d and y polymorphsdepending on heat treatment times and temperatures as shown in Table 21.1.Bulk nucleation was observed to be the dominant nucleation mechanism.The activation energy for the crystallisation process was found to be 834 KJ/mol.

Hampshire et al. (1994) studied the crystallisation behaviour of a glass ofcomposition (in e/o) 35Y:45Si:20Al containing 23 e/o nitrogen which resultedin formation of multi-phase composites, depending on the temperature ofheat treatment. B-phase (Y2SiAlO2N), Iw-phase (Y2Si3Al2(O,N)10, i.e. 10e/o N) and N-wollastonite (YSiO2N) are formed at temperatures below 1200∞C,while a-yttrium disilicate (Y2Si2O7), N-apatite (Y5Si3O12N) and YAG(Y3Al2O12) are formed at higher temperatures. At relatively low heat treatmenttemperatures of ~950–1100∞C, the nucleation and growth of N-wollastoniteand the intermediate phases B and Iw are kinetically favoured over the morestable equilibrium phases YAG and Si2N2O. In a study of the initial stages ofcrystallisation of this same glass (Besson et al., 1997), it was found that thecreep rate is higher than for the parent glass, since the residual glass in thecomposite following nucleation has a lower viscosity due to a decrease inyttrium and an increase in impurity cations. Following complete crystallisation,the creep rate was very low, showing that very little glass remains in thecomposite.

Further studies on the crystallisation of B and Iw phase composites fromY-Si-Al-O-N (and Er-Si-Al-O-N) glasses with similar compositions to those

Table 21.1 Crystalline phases observed in 28Y:56Si:16Al:83O:17N glass-ceramic composite after two-stage heat treatment process

HT temperature Crystalline phases(Tg = 985∞C)

Tg a-Y2Si2O7, b-Y2Si2O7, Si2N2OTg + 20 K a-Y2Si2O7, b-Y2Si2O7, Si2N2O, YAG*/AlYO3

*

Tg + 40 K b-Y2Si2O7, a-Y2Si2O7, Si2N2O, YAG*/AlYO3*

Tg + 60 K b-Y2Si2O7, a-Y2Si2O7, Si2N2O, YAG*/AlYO3*

Tg + 80 K b-Y2Si2O7, Si2N2O, g-Y2Si2O7, YAG*/AlYO3*

Tg + 100 K b-Y2Si2O7, Si2N2O, g-Y2Si2O7 YAG*/AlYO3*

*Trace amounts.


reported above have been undertaken (Lemercier et al., 1997; Young et al.,2000; MacLaren et al., 2001; Diaz and Hampshire, 2002; Díaz et al., 2003).Glasses with 5–20 e/o N heat-treated in the range 1000–1250∞C showdevelopment of different phase assemblages depending on N content,temperature and modifying cation. As shown in Fig. 21.3 (Menke andHampshire, 2005), at low nitrogen contents (up to 10 e/o), Iw-phasepredominates. At 1050∞C and a nitrogen content of 12 e/o, B-phase appearsin addition to Iw-phase. A further increase in nitrogen content diminishes theIw-phase, and B-phase predominates, becoming the only phase in the glass-ceramic at higher nitrogen contents (17–20 e/o N).

In the Y-Si-Al-O-N system, although B-phase is the major phase at 1050∞C,at 14 to 17 e/o N, Iw-phase predominates at temperatures above 1100∞C andtraces of b-Y2Si2O7 are observed above 1150∞C. At 20 e/o N, B-phasepredominates at 1050–1100∞C but is balanced by formation of Iw and otherphases such as N-apatite at higher temperatures. At crystallisation temperatureshigh enough to nucleate the stable equilibrium phases (~1250∞C), thecrystallisation process involves a partitioning of the glass into oxide phases(Y2Si2O7, Y3Al5O12, Al6Si2O13 and Al2O3) and a high nitrogen phase silicon

B-phase (002)

20 e/o N

17 e/o N

lw-phase (020)

14 e/o N

12 e/o N

10 e/o N

8 e/o N

Inte

nsi

ty

15 16 17 18 19 20 212q (∞)

21.3 XRD traces of the 35Y:45Si:20Al:xO:yN glass-ceramics heattreated at 1050∞C showing the major peaks for the B and Iw phases.


oxynitride, with its structure made up of SiON3 tetrahedra. At low crystallisationtemperatures (~1100∞C), this partitioning is inhibited by the much loweratomic diffusion rates, and the crystallisation process involves the nucleationand growth of phases that are kinetically preferred owing to their compositionalsimilarity to structural units shown to be already present in the parent glass.

Properties of Y-Si-Al-O-N glass-ceramic composites

Crystallisation of Y-Si-Al-O-N glasses results in glass-ceramic compositeswhich will offer greater refractoriness and better mechanical properties thanthe starting glass. As shown above, the crystallisation of these glasses iscomplex, and different composite products form depending on initial glasscomposition and on the temperature of the heat treatments (Thomas et al.,1982; Leng-Ward and Lewis, 1985; Dinger et al., 1988; Besson et al.,1993;Hampshire et al., 1994; Hampshire, 1994; Menke et al., 2005).

Properties of glass-ceramic composites with cation composition (in e/o)35M = Y or Er:45Si:20Al have been measured (Menke and Hampshire,2005). A multiphase Y-Si-Al-O-N (14 e/o N) glass-ceramic composite of Iwand YAG formed by heat-treatment at 1200∞C exhibits very good mechanicalproperties with Young’s modulus of 188 GPa. An equivalent composition Er-Si-Al-O-N (20 e/o N) glass-ceramic composite containing ErAG, N-wollastonite and N-apatite exhibits a Young’s modulus of 204 GPa.

Besson et al. (1993) found that, in general, the thermal expansion coefficientof the Y-Si-Al-O-N glass-ceramic composite containing Y2Si2O7, YAG andSi2N2O varies linearly with temperature between 20 and 800∞C and is similarto that of the parent glass. At higher temperatures the thermal expansioncoefficient increases slightly. The hardness of Y-Si-Al-O-N glass-ceramiccomposites was 9–10 GPa. This is quite high for a glass-ceramic and may becompared with the hardness values for b-sialon (15 GPa) and SiC (25 GPa).Unlike many glass-ceramics, the composite exhibited a positive temperaturecoefficient of electrical resistivity.

The fracture behaviour of glasses of composition (in e/o) 35Y:45Si:20Alwith 17 e/o N was assessed by means of flexural strength measurements andusing a fractographic approach (Hampshire, 2003b). A first-stage heat treatmentat Tg + 20∞C (960∞C) carried out on polished flexural specimens, in order toround out the crack tips produced during the surface-finishing step, resultedin substantial improvements in strength. A two-stage glass-ceramic heattreatment (960∞C for 1 hour and 1050∞C for 5 hours), which allows crystalgrowth to form the B-phase, was carried out and a mean flexural strength of549 MPa was measured for the glass-ceramic.


21.3.4 Glass-ceramic composites in Ln-Si-O-N andLn-Si-Al-O-N systems

Crystallisation to form composites in the Ln-Si-O-N and Ln-Si-Al-O-N systemshas been reported. The disilicate, metasilicate, N-apatite and N-wollastonitephases found in the Y-Si-O-N system (see Fig. 21.2) also exist in the Ln-Si-O-N systems (Leng-Ward and Lewis, 1990; Mandal et al., 1992a; Sun et al.,1995; Liddell et al., 1998).

The melilite phases, of compositions Ln2Si3O3N4, occur across the entirerange of rare earths with the exception of lanthanum (Mandal et al., 1992a).J-phases, of composition Ln4Si2O7N2, are similar, but the La-J-phase doesoccur even though it has a limited stability (Mandal et al., 1992a).

N-apatites (Y5Si3O12N) occur across the whole series of lanthanides; atthe low atomic number (low Z) end, there is a triangular solid solution rangeextending between Ln8(SiO4)6, Ln9.33(SiO4)6O2 and Ln10(SiO4)6N2 (Morrisseyet al., 1990; Mandal et al., 1992a; Hampshire et al., 1992). The N-wollastonitesof composition LnSiO2N are the least stable of the Ln-Si-O-N oxynitrides,extending from La only as far as Sm, but the La, Ce and Nd end-membershave much better thermal stability than the yttrium N-wollastonite (Mandalet al., 1992a). The higher-Z rare-earth cations (La, Ce, Nd) form compoundsoriginally thought to be of composition Ln2O3.2Si3N4, but more recentlyshown to be of composition Ln3Si8N11O4 (Mandal et al., 1992b); analoguesof these do not occur in the Y-Si-Al-O-N system.

The extension of all these phases into the five-component Jänecke prismby simultaneous substitution of Si by Al and of N by O has been reported,especially for the wollastonite series. Korgul and Thompson (1989) exploredcrystallisation of the end-member LaSiO2N, CeSiO2N and NdSiO2Nwollastonites in considerable detail. Work has also been reported on thesialon U-phases (La3Si3–xAl3+xO12+xN2–x) (Spacie et al., 1988; Leng-Wardand Lewis, 1990; Mandal et al., 1992a; Ramesh et al., 1996) and the sialonW-phases (Mandal et al., 1992a). During crystallisation of Ln-Si-Al-O-Nglasses to form glass-ceramic composites, for larger radius cations (La, Nd,Sm, etc.) the W-phase (Ln4Si9Al5O30N) forms, whilst for smaller radiuscations (Y, Er, etc.) the B-phase (Ln2SiAlO5N) or disilicates (Ln2Si2O7) aremore stable.

With respect to disilicates, Liddell and Thompson (1986) evaluated theeffects of cation radius on the stability of various yttrium and lanthanidedisilicates and endorsed earlier work which shows that, in general, a-polymorphs are stable for large radius cations whilst b-polymorphs are stablefor small radius cations. Properties of some Ln-Si-Al-O-N glass-ceramiccomposites (of original composition 28Ln:56Si:16Al:83O:17N; Ln = Ho,Er, Yb, Y) after heat treatment at 1200∞C for 5 hours to form b-Ln2Si2O7 andresidual glass are compared with those for the parent glasses in Table 21.2.


Crystallisation of the Ln2Si2O7 results in substantial increases in elasticmoduli (tensile and shear), with Yb and Y-Si-Al-O-N glass-ceramics exhibitingYoung’s moduli of 175 GPa.

For M1-M2-Si-Al-O-N glasses where M1 = La or Nd and M2 = Y or Er,temperatures at which crystallisation exotherms arise have also been determined(Pomeroy et al., 2005) as well as crystalline phases present after the glasseshad been heat treated to 1300∞C in nitrogen. The results clearly demonstratethat glass properties vary linearly with effective cation field strength of thecombined modifiers (M1, M2) which is calculated from the atomic fractionsof M1 and M2 and their associated cation field strengths. Glass stability inboth the La–Y and La–Er systems reaches a maximum at M1 and M2 fractionsof 0.5 because of the relative stability of different oxynitride and disilicatephases with changes in ionic radius. Furthermore, La appears to stabilise thea-polymorph of yttrium disilicate because of combined La-Y ionic radiuseffects. Studies conducted by Weldon et al. (1996) showed that, for mixedmodifier (La:Er = 1:1) glass, devitrification to apatite occurred rather than atwo-phase mixture of La–W phase and yttrium disilicate which were theprimary devitrification products of the single modifier (La or Er) Si-Al-O-Nglasses. Chen et al. (1997), in a study of a mixed modifer La-Y-Si-O-Nglass, also observed the crystallisation of apatite.

21.3.5 Glass-ceramic composites in theNd-Mg-Si-O-N system

Crystallisation of glasses in the Nd-Mg-Si-O-N system to form glass-ceramiccomposites has been investigated. Morrissey et al. (1990) showed that heattreatment at a single temperature resulted in only a small increase in hardnessfor a 12:24:64 Nd:Mg:Si composition, but two-stage heat treatments resultedin a much higher increase. They found that the optimum nucleation temperaturewas related to the glass transition temperature of the materials (usually ~ Tg

Table 21.2 Comparison of density, r, microhardness, Hv, Young’s modulus, E, andshear modulus, G, for LnSiAlON glasses and glass-ceramic composites

Glass- Before heat treatment After heat treatmentceramic (1200∞C for 5 h)composite ————————————— ————————————————————system r Hv E G r Hv E G Crystalline

(g cm–3) (GPa) (GPa) (GPa) (g cm–3) (GPa) (GPa) (GPa) phasespresent

HoSiAlON 5.05 9.67 145 56 5.06 9.95 166 65 Ho2Si2O7

ErSiAlON 5.10 10.2 146 57 5.04 10.5 162 59 Er2Si2O7

YbSiAlON 5.15 10.15 142 55 5.20 10.4 175 68 Yb2Si2O7

YSiAlON 3.73 10.13 145 56 3.78 10.25 175 69 Y2Si2O7


+ 40 K). The major phases formed included apatite, which can form a rangeof solid solution from an oxide form containing Mg to the N-apatite,Nd5Si3O12N.

Lonergan et al. (1992) optimised the heat-treatment schedule of Nd-Mg-Si-O-N glass-ceramic composites with a cationic composition in equivalentpercent of 36 (Nd+Mg):64Si and different nitrogen contents. They found aswell that a two-stage heat treatment has to be applied to optimise themorphology of phases in the composite and therein the mechanical properties.The nucleation temperature was related to the glass transition temperature.The crystallisation temperatures increased initially with nitrogen content butthere seemed to be a levelling off at higher nitrogen contents. These glasseswere also used as matrices for glass–SiC composites as described in the nextsection.

21.4 Oxynitride glass–silicon carbide composites

Friend et al. (1990) were the first to report on oxynitride glass and glass-ceramic composites reinforced with SiC whiskers and showed that Young’smodulus increases in line with the volume fraction of SiC additions. However,at the temperatures used for hot-pressing, interaction of SiC with the oxynitrideglass (represented by SiO2) occurs according to the following:

6SiC + 3SiO2 + 6N2 = 3Si3N4 + 6CO (21.3)

and with a higher glass:SiC ratio:

3SiC + 3SiO2 + 2N2 = Si3N4 + 3CO + 3SiO (21.4)

The fracture toughness increased in line with increases in Young’s modulus,E, but not as a consequence of increases in fracture surface energy, gf.

Hampshire and Ramesh (1996) and Schneider et al. (1999) studied glassesand glass-ceramics in the Y-Si-Al-O-N-C and Ln-Si-Al-O-N-C systems. Table21.3 shows the compositions investigated to form glasses by melting at1700∞C under nitrogen atmosphere as for previously reported oxynitrideglasses and the phases crystallised after heat treatments at 1200∞C for 1 h.The major phases formed were various polymorphs of yttrium disilicate(principally y- and b-Y2Si2O7) and mullite (3Al2O3.2SiO2). The maximumsolubility of carbon in these glasses is less than 5 e/o, so no carbide phaseswere apparent except in the composition containing 10 e/o C. The compositioncontaining 10 e/o N had not crystallised, showing that nitrogen improvesglass stability.

Rouxel and Verdier (1996) studied the viscoplastic forming and thecrystallization ranges of SiC particle reinforced Y-Mg-Si-Al-O-N glasscomposites produced by hot-pressing at ~1050∞C. Crystallization starts beyond1050∞C, with spinel, MgAl2O4, enstatite, MgSiO3, and y- and d-Y2Si2O7 as


the major phases. A non-Newtonian flow behaviour was observed from 800to 900∞C, i.e. above and below Tg (Tg = 863∞C), with n approximately equalto 0.6 (shear thickening behaviour). The viscosity versus temperature curves(deduced from creep tests) show that the temperature must be higher than950∞C for the viscosity to be lower than 1010 Pa s. In the light of theseresults, the feasibility for viscoplastic forming of a 40 vol% SiC compositewas demonstrated by shaping a parabolic shell at 980∞C within 30 min.

Rouxel et al. (2000) reported on SiC particle reinforced Y-Mg-Si-Al-O-Noxynitride glass composites and found them to have remarkable mechanicalproperties and to be suitable for viscoplastic forming. In order to betterunderstand the complex nature of flow in these composites, the stress relaxationand creep behaviour were characterized for SiC particle sizes of 3 to 150 mmand volume fractions from 0 to 40% SiC. The viscosity coefficient, as calculatedfrom relaxation data, is very close to the creep viscosity, as determined fromthe stationary creep regime. The presence of rigid particles results in significantdecreases in the relaxation kinetics and creep rates. The smaller the particlesize or the higher the particle volume fraction, the lower the flow kineticsbecomes. Furthermore, the apparent viscosity of the composite exhibits astrain-hardening behaviour, and a critical strain, at which flow is apparentlyblocked (depending on both particle size and volume fraction), has beensuccessfully introduced to interpret the data.

Baron et al. (2000) studied the fracture behaviour of YMgSiAlON glass–SiC composites; it was assessed by means of flexural strength measurementsand a fractographic approach. The composites were produced by hot pressinga mixture of glass and SiC powder in a graphite cell lined with boron nitride,between two molybdenum sheets. The pressure was limited to 15 MPa andthe composites were hot-pressed at a sintering temperature of 1040∞C for 30minutes. Then the glass composite was annealed at a temperature of 830∞Cfor 30 minutes. Three-point bending test bars were prepared with differentpolished surfaces. Some samples were submitted to a flame polishing heattreatment. Results of flexural strength tests are shown in Fig. 21.4 (Baron etal., 2000).

Table 21.3 Phase assemblage of glass-ceramics with constant cation ratios(16.5Y:56Si:27.5Al e/o) and increasing levels of nitrogen, N, and carbon, C, after heattreatment at 1200∞C for 1 h (Hampshire and Ramesh, 1996)

Composition (eq %) Crystalline phase assemblage

16.5Y:56Si:27.5Al:100O y-Y2Si2O7, 3Al2O3.SiO2, b-Y2Si2O7, d-Y2Si2O7

16.5Y:56Si:27.5Al:98O:2C y-Y2Si2O7, b-Y2Si2O7, 3Al2O3.SiO2

16.5Y:56Si:27.5Al:98O:2N y-Y2Si2O7, b-Y2Si2O7, 3Al2O3.SiO2

16.5Y:56Si:27.5Al:90O:10C b-Y2Si2O7, y-Y2Si2O7, 3Al2O3.SiO2, Y2C3, SiC16.5Y:56Si:27.5Al:90O:5C:5N y-Y2Si2O7, 3Al2O3.SiO2, b-Y2Si2O7

16.5Y:56Si:27.5Al:90O:10N Amorphous


Fracture originated from the particles and their strength changed withsurface finish only when the defects created by machining were bigger thanthe defects introduced by the particles. The consequence is that compositesare less sensitive to machining flaws, but, when the surface finish was good,the strength of the matrix was only increased for small particle sizes. In glassmatrix particulate composites, the size of critical defects is related to the sizeof the particles leading to a decrease of strength when the particle sizeincreases as observed in Fig. 21.4. The increase in strength with increasingvolume fraction of particles is shown in Fig. 21.5 (Baron et al., 2000) and ismainly due to the resultant increase in Young’s modulus. For low particlevolume fraction, a reduction of the fracture strength was observed afterflame polishing compared to the matrix. In this case, the increase of Young’smodulus was not sufficient to compensate for the increase of critical defectsize due to the introduction of the particles. The quality of the surface finishinfluenced the fracture strength of the composites only when the flawsintroduced by the polishing became smaller than the flaws introduced by theparticles. This means that the larger the particle size the lower the fracturestrength sensitivity of the composites to machining flaws.

Matrix 3 mm 6 mm 16 mm 31 mm 150 mm

350

300

250

200

150

100

50

s r (

MP

a)

after grindingpolished 3 mmpolished 1/4 mmflame polished

21.4 Dependence of the fracture strength of the YMgSiAlON glass–SiC composites (28 vol% SiC) on the SiC particle size and the surfacefinish (after Baron et al., 2000).

21.5 Dependence of the fracture strength of the composites on thevolume fraction of particles (6 mm SiC).

polished 3 mmflame polished

Glass 10 vol% 28 vol%

s r (

MP

a)

350

300

250

200

150

100

50


21.5 Conclusion

Oxynitride glasses are silicate or alumino-silicate glasses in which oxygenatoms in the glass network are partially replaced by nitrogen atoms. Asnitrogen increases, glass transition temperature, elastic modulus, viscosityand hardness increase while thermal expansion coefficient decreases.

Many studies on crystallisation of oxynitride glasses have been carriedout. Some of these have identified suitable two-stage heat treatments fornucleation and growth of crystal phases to form glass-ceramic compositeswith significant increases in mechanical properties over the parent glass.Overall, the thermal, electrical and mechanical properties of the oxynitrideglass-ceramic composites differ from system to system and have to beestablished for every composition and every heat-treatment schedule applied.However, by using established techniques for processing of glass-ceramics,with a first-stage heat treatment to provide nucleation and a second-stageheat treatment for crystal growth, microstructures and properties of multi-phase composites can be optimised.

SiC particle reinforced oxynitride glass composites have been investigatedand found to have higher values of mechanical properties which are relatedto the volume fraction of SiC inclusions, provided that SiC–glass reactionscan be avoided. These composites have been shown to be suitable forviscoplastic forming.

21.6 References

Baron, B., Lemercier, H., Veyrac, C., Pomeroy, M., Hampshire, S. (2000), ‘Fracture ofoxynitride glasses and SiC particulate composites’, Mater. Sci. Forum, 325, 295–301.

Besson, J.L., Billiers, D., Rouxel, T., Goursat, P., Flynn, R., Hampshire, S. (1993),‘Crystallization and properties of a Si-Y-Al-O-N glass-ceramic’, J. Am. Ceram. Soc.,76, C2103–5.

Besson, J.-L., Lemercier, H., Rouxel, T., Troillard, G. (1997), ‘Yttrium sialon glasses:Nucleation and crystallisation of Y35Si45Al20O83N17’, J. Non-Cryst. Sol., 211, 1–21.

Chen, J., Wei, P., Huang, Y. (1997), ‘Formation and properties of La-Y-Si-O-N oxynitrideglasses’, J. Mater. Sci. Lett., 16, 1486–8.

Diaz, A., Hampshire, S. (2002), ‘Crystallisation of M-SiAlON glasses to Iw-phase glass-ceramics: preparation and characterization’, J. Mater. Sci., 37, 723–30.

Diaz, A., Dolekcekic, E., Pomeroy, M.J., Hampshire, S. (2003), ‘Effect of compositionand processing conditions on the Formation of Y and Er-SiAlON B and Iw PhaseGlass-ceramics’, Key Eng. Mater., 237, 247–252.

Dinger, T.R., Rai, R.S., Thomas, G. (1988), ‘Crystallization behaviour of a glass in theY2O3–SiO2–AlN system’, J. Am. Ceram. Soc., 71, 236–44.

Drew, R.A.L., Hampshire, S., Jack, K.H. (1981), ‘Nitrogen Glasses’, Proc. Brit. Ceram.Soc., 31, 119–32.

Friend, S.J., Piller, R.C., Briggs, A., Davidge, R.W., Lonergan, J.M., Hampshire, S.(1990), ‘Sintering of neodymia-containing oxynitride glasses and glass-ceramics andeffects of additions of silicon carbide whiskers’, Silicates Industriels, (Journal of theBelgium Ceramic Society) LV, 303–8.


Hampshire, S. (1993), ‘Oxynitride glasses and glass-ceramics’, in Chen I.W., Becher,P.F., Mitomo, M., Petzow, G., Yen, T-S. (eds), Silicon Nitride Ceramics – Scientificand Technological Advances, Mater. Res. Soc. Symp. Proc., 287, 93–100.

Hampshire, S. (1994), ‘Nitride ceramics’, in Swain, M.V. (ed.) Structure and Propertiesof Ceramics, Materials Science and Technology Series, Vol. 11, Weinheim, VCH,Chapter 3, 119.

Hampshire, S. (2003a), ‘Oxynitride glasses – a review’, J. Non-Cryst. Sol., 316, 64–73.Hampshire, S. (2003b), ‘SiAlON glasses, their properties and crystallisation’, Key Eng.

Mater., 237, 239–46.Hampshire, S., Pomeroy, M.J. (2004), ‘Effect of composition on viscosities of rare earth

oxynitride glasses’, J. Non-Cryst. Sol., 344, 1–7.Hampshire, S., Ramesh, R. (1996), ‘Oxynitride liquids, glasses and glass-ceramics’,

invited presentation at 1st International Workshop on Synergy Ceramics, Nagoya,Japan, November, 1996.

Hampshire, S., Drew, R.A.L., Jack, K.H. (1985), ‘Oxynitride glasses’, Phys. Chem. Glass.,26, 182–6.

Hampshire, S., Flynn, R., Lonergan, J., O’Riordan, A. (1992), ‘Oxynitride glass systemsand subsequent glass-ceramic heat treatments’, in Carlsson, R. (ed.), Ceramic Materialsand Components for Engines, Proc. 4th Int. Symp. Göteborg, June 1991, London,Elsevier Applied Science.

Hampshire, S., Nestor, E., Flynn, R., Besson, J.-L., Rouxel, T., Lemercier, H., Goursat, P.,Sebai, M., Thompson, D.P., Liddell, K. (1994), ‘Yttrium oxynitride glasses: propertiesand potential for crystallisation to glass-ceramics’, J. Eur. Ceram. Soc., 14, 261–73.

Korgul, P., Thompson, D.P. (1989), ‘The transparency of oxynitride glasses’, J. Mater.Sci., 28, 506–12.

Lemercier, H., Ramesh, R., Besson, J.-L., Liddell, K., Thompson, D.P., Hampshire, S.(1997), ‘Preparation of pure B-phase glass-ceramic in the yttrium–sialon system’, KeyEng. Mater., 132–136, 814–7.

Leng-Ward, G., Lewis, M.H. (1985), ‘Crystallisation in Y-Si-Al-O-N glasses’, J. Mater.Sci. Eng., 71, 101–11.

Leng-Ward, G., Lewis, M.H. (1990), ‘Oxynitride glasses and their glass-ceramic derivatives’,in Lewis, M.H. (ed.), Glasses and Glass-ceramics, London, Chapman & Hall.

Liddell, K., Thompson, D.P. (1986), ‘X-ray diffraction data for yttrium silicates’. Brit.Ceram. Trans. J., 85, 17–22.

Liddell, K., Thompson, D.P., Wang, P.L., Sun, W.Y., Gao, L., Yan, D.S. (1998), ‘J-phasesolid solution series in the Dy-Si-Al-O-N system’, J. Eur. Ceram. Soc., 18, 1479–92.

Lonergan, J., Morrissey, V., Hampshire, S. (1992), ‘Optimisation of heat-treatment schedulesfor oxynitride glass-ceramics’, Brit. Ceram. Proc., 49, 57–62.

MacLaren, I., Falk, L.K.L., Diaz, A., Hampshire, S. (2001), ‘Effect of composition andcrystallization temperature on microstructure of Y- and Er-SiAlON Iw-phase glass-ceramics’, J. Am. Ceram. Soc., 84, 1601–8.

Mandal, M., Thompson, D.P., Ekström, T. (1992a), ‘Heat treatment of Ln-Si-Al-O-Nglasses’, Key. Eng. Mater., 72-74, 187–203.

Mandal, M., Thompson, D.P., Ekström, T. (1992b), ‘Heat treatment of sialon ceramicsdensified with higher atomic number rare earth and mixed yttrium/rare earth oxides’,Brit. Ceram. Proc., 49, 149–62.

Menke, Y., Hampshire, S. (2005), ‘Ln-Si-Al-O-N glass-ceramics: crystallisation andproperties’, submitted to J. Eur. Ceram. Soc.


Menke, Y., Falk, L.K.L., Hampshire, S. (2005), ‘The crystallisation of Er-Si-Al-O-N B-phase glass-ceramics’, accepted by J. Mater. Sci., in press.

Morrissey, V., Lonergan, J., Pomeroy, M.J., Hampshire, S., (1990), ‘Crystallisationtreatments for neodymia-containing glasses and glass-ceramics’, Brit. Ceram. Proc.,45, 23–9.

Ohashi, M., Hampshire, S. (1991), ‘Formation of Ce-Si-O-N glasses’. J. Am. Ceram.Soc., 74, 2018–20.

Ohashi, M., Nakamura, K., Hirao, K., Kanzaki, S., Hampshire, S. (1995), ‘Formation andproperties of Ln–Si-O-N glasses’, J. Am. Ceram. Soc., 78, 71–6.

Pomeroy, M.J., Nestor, E., Ramesh, R., Hampshires (2005), ‘Properties and crystallisationof rare earth SiAlON glasses containing mixed trivalent modifiers: J. Amer. Ceram.Soc., 88(4), 875–881.

Ramesh, R., Nestor, E., Pomeroy, M.J., Hampshire, S., Liddell, K., Thompson, D.P.(1996), ‘Potential of NdSiAlON Glasses for Crystallisation to Glass-Ceramics’, J.Non-Cryst. Sol., 196, 320–5.

Ramesh, R., Nestor, E., Pomeroy, M.J., Hampshire, S. (1997), ‘Formation of Ln-Si-Al-O-N glasses and their properties’, J. Eur. Ceram. Soc., 17, 1933–9.

Ramesh, R., Nestor, E., Pomeroy, M.J., Hampshire, S. (1998), ‘Classical and DTA studiesof the glass-ceramic transformation in a YSiAlON glass’, J. Am. Ceram. Soc., 81,1285–97.

Rouxel, T., Verdier, P. (1996), ‘SiC particle reinforced oxynitride glass and glass-ceramiccomposites: crystallization and viscoplastic forming ranges’, Acta Mater., 44, 2217–25.

Rouxel, T., Sangleboeuf, J.C., Verdier, P., Laurent, Y. (2000), ‘Elasticity, stress relaxationand creep in SiC particle reinforced oxynitride glass’ Key Eng. Mater., 171, 733–40.

Sakka, S. (1995), ‘Structure, properties and application of oxynitride glasses’, J. Non-Cryst. Sol., 181, 215–24.

Schneider, N.K., Mooney, C., Baron, B., Hampshire, S. (1999), ‘Oxynitride glass compositescontaining nano-size SiC’, Brit. Ceram. Proc., 60(1), 401–2.

Spacie, C.J., Liddell, K., Thompson, D.P., (1988), ‘The U-phase in heat-treated sialonceramics’, J. Mater. Sci. Lett., 7, 95–6.

Sun, W.Y., Yan, D.S., Gao, L., Mandal, H., Liddell, K., Thompson, D.P. (1995), ‘Subsolidusphase relationship in the systems Ln2O3–Si3N4–AlN–Al2O3 (Ln = Nd, Sm)’, J. Eur.Ceram. Soc., 65, 15, 1435–8.

Thomas, G., Ahn, C., Weiss, J., (1982), ‘Characterization and crystallization of Y-Si-Al-O-N glass’, J. Am. Ceram. Soc., 65, C185–8.

Thompson, D.P. (1989), ‘Alternative grain boundary phases for heat treated Si3N4 and b-sialon ceramics’, Brit. Ceram. Proc., eds. Davidge, R.W. and Thompson, D.P., Vol.44, 1–14.

Thompson, D.P. (1992), ‘Oxynitride glasses’, in Cable, M. and Parker, J.M. (eds) HighPerformance Glasses, Glasgow, Blackie and Sons, Chapter 5.

Tredway, W.K., Risbud, S.H. (1984), ‘Influence of atmospheres and TiO2 nucleant on thecrystallisation of Mg-SiAlON glasses’, J. Mater. Sci. Lett., 4, 31–3.

Weldon, L., Pomeroy, M.J., Hampshire, S. (1996), ‘Glasses in the rare-earth sialon systems’,Key Eng. Mater., 118–119, 241–8.

Young, W., Falk, L.K.L., Lemercier, H., Peltier-Baron, V., Menke, Y., Hampshire, S.(2000), J. Non-Cryst. Sol., 270, 6–12.

575

22.1 Introduction

Functionally graded materials (FGMs) are multifunctional materials, whichcontain a spatial variation in composition and/or microstructure for the specificpurpose of controlling variations in thermal, structural or functional properties.Also in the ceramics composites field, a wide range of functionally graded(FG) ceramics are available. Hence, a possible classification of the differentclasses is made in this chapter.

Compared to homogeneous materials a lot of advantages exist. However,there would be little point in developing such a FGM material unless it couldcompete commercially with existing homogeneous ceramic componentscurrently on the market. A wide variety of processing routes are available forFGMs and it is important to choose the right processing technique for theright application. Therefore, a large part of this chapter deals with the processingof FGM, more specifically for bulk FGMs.

The gradient composition in FGMs not only results in a spatial variationin properties but will also generate residual stresses, which will affect themechanical properties. One of the potential advantages of FG components isthe positive influence of compressive residual surface stresses on the strengthand wear resistance. A correct design of the gradient for an optimal distributionof the residual stresses is therefore important, as discussed in this chapter.

Special attention will be given to structural FGM applications, where theoperating conditions are severe. More specifically, alumina/zirconia and WC/Co FGM components will be discussed.

22.2 Functionally graded ceramics concept

As a new concept in advanced materials research and development, introducedin the mid-1980s, functionally graded materials (FGMs) are defined as materialswhose spatial distribution of microstructure and/or composition is tailoredand quantitatively controlled in order to achieve an improvement in properties

22Functionally graded ceramics

G A N N É, J V L E U G E L S andO VA N D E R B I E S T,

Katholieke University Leuven, Belgium


of the final component (Mortensen and Suresh, 1995; Neubrand and Rödel,1997; Gasik, 1995). The concept of FGMs derives from the considerationthat the service conditions and required materials performance vary with thelocation in a large number of structural components. Consider for example aturbine blade, which must withstand high non-stationary heat fluxes andcentrifugal acceleration. An ideal structure for this application would consistof a tough metal core and a corrosion resistant ceramic at the hot surface ofthe blade. If the ceramic is directly bonded to the metal, spalling may occurduring thermal cycling because of the high thermal stresses at the interface.A graded material with a smooth transition from the ceramic surface to themetal core can avoid the thermo-mechanical stress concentration at the interface,and thus has better performance. There are many other applications wherethe requirement for properties cannot be attained by a single material (Neubrandand Rödel, 1997). In all these cases, graded materials offer the possibility tocombine two materials properties, avoiding most of the disadvantages of abi-material. Fundamental and applied research in the recent past has clearlyshown that the introduction of compositional step and profile gradients inFGMs can be a benefit on many accounts (Suresh and Mortensen, 1997;Munz et al., 1998; Bao and Cai, 1997).

In a broad sense, the concept of FGMs is not new. In China, an ancientproverb is used frequently in people’s daily tasks: ‘the best steel only forcutting edges’. If we eliminate its philosophical sense, this proverb fromancient blacksmiths expressed clearly the idea of FGMs. In fact, this idea hasbeen exploited for hundreds of years in iron and steel production (Gasik,1995). Although it is not difficult to find examples throughout recordedhistory of using the idea of graded materials, as a systematic engineeringapproach it is far from being the norm. Instead, current engineering practicegenerally involves the design of a part using one of the large number ofgenerally mass-produced uniform materials. When a structure requires vastlydifferent materials, different uniform materials are joined along sharpboundaries using a variety of joining or coating methods. The task of thematerials engineer is defined as seeking to improve materials, while the taskof the structural part designer is using the handbooks of available materials.With the innovative FGM concept, component design and fabrication arebased not on a list of existing materials but on a choice of available basicmaterial ingredients and material processes, combined with three-dimensionalmechanical analysis of graded structures. These two engineering disciplinesare combined to synergistically design both the component and its processing.Here lies a second definition of the functionally graded materials concept, asan approach in engineering rather than a physical entity (Mortensen andSuresh, 1995).

Functionally graded ceramics 577

22.3 Classification of FG ceramics

There are a number of possibilities to classify graded materials (Neubrandand Rödel, 1997):

∑ According to the material classes which the graded component combines,e.g. metal/ceramic, polymer/ceramic, metal/metal, ceramic/ceramic, etc.

∑ According to the relative extension of the gradient, i.e. to what extent thegradient is distributed across a component. In functionally gradient coatingsand joints, the gradient extends only over a part of the component close toits surface or in the interior. In functionally graded bulk materials, thegradient comprises the entire part.

∑ According to the function in a component. Figure 22.1 gives some examplesof applications of graded materials, which allow unusual combinations ofproperties.

22.4 Processing of FGMs

One characteristic of the fabrication of FGMs is certainly the very widevariety of available processing methods. Functionally graded materials includematerials with a gradient in composition, grain size and/or porosity. Thegeneral goal of processing of FGMs is to realize a spatial distribution in themicrostructure and/or composition in the final part.

When selecting the processing method, differences between the propertiesof the two constituent phases of the FGM are of primary importance. In acompositional FGM, for example, the difference in heat resistance between

Tool

Turbineblades

Substrate

Sensors

Actuators

HardnessToughness

Corrosionresistance

Thermalconductivity

Electricalinsulation

Dielectricproperties

Piezoelectricproperties

Acousticimpedance

22.1 Examples of applications of graded materials (Neubrand andRödel, 1997).


the two phases is a key factor. If the two phases have a significantly differentmelting point, as in ceramic/metal FGMs, the composition gradient can beformed by producing a porosity gradient preform of the refractory phasesubsequently infiltrated by the molten second phase to get a dense finalproduct. If the two phases have a similar melting point, infiltration cannot beused because the skeleton cannot keep its strength during infiltration. Thedimensions and geometry of the FGM has to be considered as well. It isfeasible to produce FGMs in many systems with thermal coating technologies,but their low efficiency makes them useless for the production of three-dimensional bulk FGMs.

22.4.1 Shaping and consolidation processes for FGMs

For the fabrication of bulk FGMs, powder metallurgical processing is mosteconomic and suitable for mass production. In order to produce a FGM byconventional powder processing, a green body with the desired gradient inphase volume fraction is first fabricated. After shaping and consolidation,this green body has to be densified by sintering.

The gradation methods can be divided into two groups: dry and wetprocesses (Fig. 22.2) (Neubrand and Rödel, 1997). Dry processes are fast,but in general only allow the generation of step-graded profiles. In wetprocessing, a drying step is required for the removal of the liquid but continuousmixing is facilitated and smoother continuous gradients may be produced.Furthermore, transport processes occurring in suspensions, e.g. sedimentationand electrophoresis, may be used to produce gradients at low cost. The mainchallenge associated with powder processing is frequently the densification

Powder metallurgyDry processing Wet processes

Sintering

Dry

Preparationof slip

Preparation of suspension

Sequentialcasting

Spray Centrifugation Sedimentation EPD

Gravitysedimentation

Centrifugalcasting

Wetspraying

Slip casting EPDConventional

PMCentrifugal

PM

Preparationof

mixtures

Continuousmixing ofpowders

Sequentialstacking

Continuousstacking

Predensification Debinding

22.2 Powder processing routes for FGMs.


of the graded powder compact. Sintering rates differ with position and unevenshrinkage may lead to warping and cracking, unless sophisticated sinteringtechniques are employed.

A widely used technique for ceramic/ceramic gradient materials is sequentialslip casting where slips of different compositions are cast one after and overanother (Requenna et al., 1993). By using a premixing system, the castingcomposition can be tailored continuously (Chu et al., 1993).

In a process called wet spraying (Schindler et al., 1998), suspensions oftwo powders are created, mixed and sprayed under computer control on aheated substrate. After forming, the green body is removed from the substrate,for bulk FGMs, or bonded with the substrate, for FGM foils. A through-thickness composition gradient can be created by controlling the ratio of thetwo powders in the mixed suspension. Centrifugal casting (Fukui et al.,1994; Watanabe et al., 1998) is another FGM consolidation method usingsuspension mixing to realize the gradient. When suspensions of two powdersof different density or different grain size are mixed and injected into acylindrical cavity, which is rotating at high speed, the centrifugal forcescause a compositional or porosity gradient in the growing powder compactin the radial direction. Before stopping the rotation, wax is injected into thesystem to bind the powders in order to increase the green strength for bodyhandling. The porous FGMs with a gradient distribution in porosity can beused as a preform for filters, or for ceramic membranes. A process similar tocentrifugal casting is gravitational sedimentation (Bernhardt, 1999). Centrifugalcasting can only be used for cylindrically shaped parts, whereas gravitationalsedimentation is suitable for flat FGMs.

Among the different colloidal processing techniques, electrophoreticdeposition (EPD) is a very promising method (Put et al., 2003a, 2003c,2002; Vleugels et al., 2003; Anné et al., 2004) because it is a fairly rapid,low-cost process for the fabrication of ceramic coatings, monoliths, composites,laminates and functionally graded materials varying in thickness from a fewnanometres to centimetres (Van der Biest and Vandeperre, 1999).Electrophoretic deposition is a two-step process (Fig. 22.3). In the first step,particles having acquired an electric charge in the liquid in which they aresuspended are forced to move towards one of the electrodes by applying anelectric field to the suspension (electrophoresis). In the second step (deposition),the particles collect at one of the electrodes and form a coherent deposit onit. The deposit takes the shape imposed by this electrode. After drying andremoval from the electrode, a shaped green ceramic body is obtained. Firingthis green body then results in a ceramic component. Gradient materials canbe obtained since the composition of the next powder layer that deposits isdetermined by the composition of the suspension at that moment (Fig. 22.3).Judiciously adapting the powder concentration in the suspension allows oneto generate a well-controlled gradient profile in a continuous shaping step.


The process is not material specific, since a wide variety of materials havealready been deposited such as metal powders, ceramics, glasses, and polymers(Van der Biest and Vandeperre, 1999). In general, the only shape limitationis the feasibility to remove the deposit from the electrode after deposition.Continuously graded materials in the Al2O3/ZrO2 (Vleugels et al., 2003),ZrO2/WC (Put et al., 2002), and WC/Co (Put et al., 2001) system havealready been explored by means of EPD.

A prerequisite for successful production of FGM materials by means ofEPD is a full control of the kinetics of the process. Kinetic models havetherefore been developed for processing an FGM in a multi-component systemby means of EPD (Put et al., 2003b). As an example, the composition profile(Fig. 22.4) and microstructure (Fig. 22.5) of an Al2O3/ZrO2 FG (Vleugels etal., 2003) disk with a homogeneous composite core (75 vol% Al2O3), a pureAl2O3 surface layer on one side and a homogeneous composite (90 vol%Al2O3) on the other side, and intermediate symmetrically profiled gradedlayers, is presented. As will be explained below, a convex alumina gradientprofile is suggested to give the highest compressive stress in the aluminaouter layers and the lowest tensile stresses in the core of the disk. The ZrO2

(white) and Al2O3 (grey) phases can be clearly differentiated in themicrostructure. The ZrO2 phase is well dispersed in the Al2O3 matrix in thegradient parts and in the core of the FGM.

Scale: nanometres to millimetres

Gradient thickness

Gradient profile

Composition100%

EPD set-up for FGM

22.3 Electrophoretic deposition process for FGM materials.


22.4.2 Densification of FGM powder compacts

A major challenge is the densification of the graded powder compacts. Theprocessing of FGM materials by powder metallurgy methods often facedundesirable excessive bending or warping of the component after sintering(Miyamoto et al., 1999). Due to excessive thermal residual stresses, cracksand other defects may often be observed in the final FGM component unlessproperly manufactured.

Measured profilePredicted profile

0 1 2 3 4 5Sintered distance (mm)

100

95

90

85

80

75

70

65

60

Co

nte

nt

of

Al 2

O3

(vo

l%)

22.4 Measured and predicted FGM profile of an Al2O3/ZrO2 FGM disk.

1 mm

A B C D E

2mm

A B C D E

22.5 General overview and some detailed micrographs of specificlocations in the FGM disk.


Figure 22.6 shows typical cracks observed in symmetrically graded Al2O3/ZrO2 disks, shaped by electrophoretic deposition and densified by pressurelesssintering. Transverse cracks were observed in the ZrO2-rich core of thesintered symmetric TZP/Al2O3 disks. The crack propagation, however, stoppedin the outer pure Al2O3 layer, indicating that the in-plane tensile stress islocated in the centre of the disks, which should be lowered. Hillman et al.(1996) observed similar defects in symmetrical laminates with Al2O3/ZrO2

layers at the surface and a ZrO2 central layer. Cai et al. (1997a, 1997b)discussed the embryonic stage of a sinter crack as observed in Fig. 22.6(b),and found regions of low density and cavitational defects in Al2O3/ZrO2/Al2O3 laminates. These defects are most susceptible to residual tensile stressesduring cooling in the core, due to the higher coefficient of thermal expansion(CTE) of zirconia. These regions of lower density (pores) must have formedas a result of the tensile stress that develops during the differential shrinkageduring densification between the Al2O3 and the Al2O3/ZrO2 layers. The poresthen act as pre-existing flaws for the generation of thermal expansion mismatchcracks during cooling via linkage of the pores and cavitational defects.

Elimination of the transverse cracks can be accomplished by decreasingthe overall shrinkage of the composites. This is done either by decreasing thecompositional difference between the different layers (Cai et al., 1997a,1997b) or by adjusting the green density of the different layers (Novak andBeranic 2005). Another possibility is to decrease the cooling and heating rateduring sintering (Cai et al., 1997b). The mismatch stresses during the heatingcycle are decreased by the viscous nature of the FGM material at the sinteringtemperature. The sintering mismatch stress is proportional to the mismatchsintering rate. Reduced cracking under a slow cooling rate is probably due tothe relaxation of residual stresses during the initial period of cooling.

Almost all ceramic/ceramic bulk FGMs are sintered by conventionalpressureless sintering (Wu et al., 1996; Marple and Boulanger, 1994; Cichockiand Trumble, 1998) or hot pressing (Kawai and Wakamatsu, 1995; Vanmeensel

(a) (b)

Crack

Al2O3

Gradient

Al2O3/ZrO2

Homogeneous

Gradient

Homogeneous

Gradient

Homogeneous

22.6 (a) FGM schematic; (b) typical crack observed in Al2O3/TZP FGM.


et al., 2004), depending on the sintering properties of the two components.In metal/ceramic FGMs with a continuous metal phase and a discontinuousceramic phase, the sintering rates are controlled by the densification of themetal phase and such FGMs can be densified by conventional sinteringmethods (Neubrand and Rödel, 1997). In most FGMs where a high ceramicphase content is envisaged, however, some special approaches have to beconsidered for full densification.

In addition to conventional sintering, reactive powder processing, alsocalled combustion synthesis or self-propagating high-temperature synthesis(SHS), can be used if the target compounds can be synthezised from thestarting powder mixture (Stangle and Miyamoto, 1995). This process comprisesa rapid and exothermic chemical reaction to simultaneously synthesize someor all of the constituent phases in the FGM and densify the component.

A more advanced technique, such as Spark Plasma Sintering (SPS) orPulsed Electric Current Sintering (PECS) (Tokita, 1999), is also used forFGM fabrication. It is a pressure-assisted sintering method in which a highcurrent is pulsed through a die/punch/sample set-up, which can be comparedwith that of conventional hot pressing. The large current pulses generatespark plasmas, a spark impact pressure and Joule heating. The sinteringmechanism and mechanical properties of the sintered compacts showcharacteristics different from conventional pressure-assisted sintering processesand parts. This technique offers significant advantages for various kinds ofnew materials and consistently produces a dense compact in a shorter sinteringtime and with finer grain size than conventional methods. Spark plasmasintering of FGMs uses a temperature gradient in the system, which allowsa homogeneous densification of FGMs by matching the temperature gradientto the shrinkage rate gradient of the compact. With a spark plasma system,large ceramic/metal bulk FGMs (~100 mm across) can be homogeneouslydensified in a short time with heating and holding times totalling less thanone hour. Amongst the reported spark plasma sintered systems are WC-based materials (WC/Co, WC/Co/steel, WC/Mo), ZrO2-based composites(ZrO2/steel, ZrO2/TiAl, ZrO2/Ni), Al2O3/TiAl, etc. (Tokita, 1999). A systematicintroduction to spark plasma sintering can be found in a review by Tokita(1999) or a recent paper by Hennicke and Kessel (2004), where the process,mechanical properties, size and shape effects, and production machine systemsare reported.

Microwave sintering is another promising technique for ceramic/metalFGMs to eliminate the difficulty of inequality of the shrinkage rate. As anewly developed sintering technique, microwave sintering uses microwave+irradiation to heat the ceramic or ceramic-based composite compact (Gerdesand Willert-Poradu, 1994; Willert-Porada, 1999; Zhao et al., 2000). Themechanism of microwave heating is based on the dielectric loss of the ceramicphases involved, resulting in a volumetric heating technique in which theheat is generated by the compact itself.


22.5 FGM design for structural applications

FGMs for ceramics offer great promise in applications where the operatingconditions are severe, for example wear-resistant linings for handling largeheavy abrasive ore particles, high-speed cutting tools, rocket heat shields,heat exchanger tubes, thermoelectric generators, heat-engine components,plasma facings for fusion reactors, and electrically insulating metal/ceramicjoints. They are ideal for minimizing the thermo-mechanical mismatch inmetal–ceramic bonding. As presented in Fig. 22.7 (Munz et al. 1998), astress discontinuity is generated in case of a bi-material combination, whichmay result in a poor adhesion or even delamination, whereas a significantlylower stress concentration can be established for a graded interface, therebyimproving the materials adhesion and reducing the probability of delamination.

The residual stresses have a large influence on the properties of a structuralbody, e.g., a correct design of the composition gradient can generatecompressive stresses at those locations which are loaded in tension duringapplication. Compressive stresses at the surface can also have a beneficialeffect on the tribological properties of the component (Novak et al., 2005).

22.5.1 Thermal stress calculations in FGM material

The calculation of thermal stresses in functionally graded materials is alreadya relatively old topic (Yang et al., 2003). Two methods can be distinguished:analytical methods and finite element methods. However, the complicatingeffect of the elastic modulus variation with the position severely limits thescope of problems that can be solved analytically. Therefore, the majority ofthe analytical work has been for FGM films or other simple structures (Beckeret al., 2000). Analytical models have been developed for the calculation ofthermal stresses for 1-D FGM symmetrical plates (Jung et al., 2003), non-

–0.300 –0.200 –0.100Y (mm)

120

100

80

60

40

20

0

s y (

MP

a)

–0.300 –0.200 –0.100Y (mm)

200

150

100

50

0

s y (

MP

a)

22.7 Stress distribution in (a) a coating–substrate system and (b) aFGM (Munz et al., 1998).


symmetrical plates (Ravichandran, 1995), cylinders (Chen and Awaji, 2003)and spheres (Obata and Noda 1994). Finot and Suresh (n.d.) have written asoftware program, called Multitherm, to calculate thermal stresses in a planestress, plane strain or a biaxial state.

For a more general 2-D or 3-D problem, numerical methods like finiteelement analysis are required. These are, costly however, since a full analysisfor each material pair, geometry and gradient must be performed.

In this section, an analytical solution to calculate residual stresses in anFGM disk is discussed, based on simple linear elastic plate theories of classicalmechanics, and used for the calculation of residual stresses in a plane stressstate. An equi-biaxial stress analysis differs from a plane stress state bysimply replacing the Young’s modulus E by the corresponding biaxial modulusE¢ = E/(1 – n). In this way, the residual thermal stress can be calculated in thecentre of the FGM disk, far enough away from the free edges where acomplex stress state is present.

Consider therefore an initially perfectly planar unconstrained, layered orgraded plate as presented in Fig. 22.8, which is subjected to a uniformthermal excursion. During cooling from the sintering temperature, the layerswith the highest CTE will contract more than the layers with a lower CTE.This causes a variation in thermal stress along the thickness of the plate. Theplate can bend or warp due to the through-thickness strain gradient, therebyaccommodating the thermal stress. When the geometrical conditions of theplate (isotropic in-plane) are such that the strain is allowed to be a functionof z only, the in-plane normal strain can be expressed as:

e (z) = exx = e0 + bz (22.1)

where e0 is the normal strain at z = 0, and b is the curvature of the plate in

y

x

z

h1

h2

E1, n1, a1

E2, n2, a2

E3, n3, a3

E3, n3, a3

E2, n2, a2

E1, n1, a1

22.8 Schematic of a layered disk.


its plane. The stress s(z) = sxx = syy for the biaxial stress state is given by(Giannakopolous et al., 1995; Suresh et al., 1994):

s n e a( ) = ( )

1 – ( ) [ ( ) – ( ) ( )]z

E zz

z z T zD

= ( )

1 – ( )[ + – ( ) ( )]0

E zz

z z T zn e b a D (22.2)

where a is the coefficient of thermal expansion (CTE), DT is the differencebetween the sintering temperature and room temperature, n is the Poissoncoefficient, and E is the Young’s modulus.

The resultant force and the resultant moment of the stress distributions(z) along the height z must be equal to the applied axial force, Fap, and theapplied bending moment, Map, respectively:

–

–ap

2

1

( )d = h

h

z z FÚ s (22.3)

–

–ap

2

1

( ) d = h

h

z z z MÚ s (22.4)

These two conditions, i.e. force balance and moment balance, lead to a linearsystem of equations for e0 and b. The solutions, in the absence of buckling,are given by (Giannakopolous et al., 1995; Suresh et al., 1994):

e 02 0

ap1 1

ap

12

0 2 =

– ( + ) + ( + ) –

I M F I M MI I I

(22.5)

b = ( + ) – ( + )

– 1 0

ap0 1

ap

12

0 2

I M F I M MI I I

(22.6)

with:

I z E z zih

hi = ( )d

– 1

2

Ú for i = 0, 1, 2 (22.7)

M z z T z E z zih

hi = ( ) ( ) ( )d

– 1

2

Ú a D for i = 0, 1 (22.8)

From these equations, the thermal stress profile through the thickness ofthe multi-layered material with a compositional profile can be calculated. IfE and aDT vary continuously with z, s is continuous too. Because, theeffective properties change as functions of the position in an FGM material,it is necessary to estimate the local composite properties as functions of


composition and temperature. Rules of mixtures can be used to estimatethese properties (Suresh and Mortensen, 1998).

This analytical model is validated by calculating the stress profile in anAl2O3/ZrO2 FGM disk with homogeneous pure Al2O3 surface layers, ahomogeneous Al2O3/ZrO2 composite core and intermediate graded layerswith a convex, linear and concave profile. The material properties used inthese calculations are given in Table 22.1. The composition profiles andcorresponding residual stresses for the FGM disks sintered at 1550∞C aregiven in Fig. 22.9. These calculations reveal that the FGM disks with aconvex (n = 0.5) alumina gradient profile result in the highest compressivestress in the alumina outer layers and the lowest tensile stresses in their core.

22.5.2 Al2O3 and ZrO2-based ceramic/ceramic FGMs

ZrO2-based FGMs have been of high interest (Marple and Boulanger, 1994,Anné et al., 2004, Zhao et al., 2000, Put et al., 2002) because of the hightoughness and excellent strength of tetragonal ZrO2 polycrystalline materials(TZP). ZrO2-based FGMs are developed mainly for energy conversion systems,biomedical applications and cutting tools, where high hardness has to becombined with high toughness or where thermal stresses have to be released(Sanchez-Herencia et al., 2000).

Al2O3/ZrO2 FGMs were studied intensively for biomedical applications,

Table 22.1 Material properties of Al2O3 and ZrO2

Young’s modulus, E Poisson’s ratio, n CTE, a(GPa) (10–6 K–1)

Al2O3 360 0.237 7.40ZrO2 180 0.300 1.01

0 2 4 6

Distance (mm)(b)

n = 1.75n = 1n = 0.5

150

100

50

0

–50

–100

–150

Res

idu

al s

tres

s (M

Pa)

n = 1.75n = 1n = 0.5

0 1 2 3 4 5Distance x (mm)

(a)

100

95

90

85

80

75

Al 2

O3

con

ten

t (v

ol%

)

22.9 Composition and corresponding residual stresses for a FGM diskwith a convex (n = 0.5), linear (n =1) and concave (n = 1.75)composition profile.


specifically for implants like hip prostheses (Anné et al., 2004). The FGMidea was applied to increase the performance of total hip replacement prosthesesthrough design and development of alumina/zirconia functionally gradedmaterials using electrophoretic deposition as a near-net shaping technique(Fig. 22.10). Today, most of the joint prostheses consist of metallic componentsarticulating against polymer counterparts, of which more than 10% needrevision after only 10 years. This is expensive and reduces the quality of lifeof the patients. Also the clinical effects of metallic ion-releasing implantsand polymer wear debris are a matter of concern. This demonstrates the needfor more wear-resistant and biocompatible materials such as ceramics. Thereforealumina–zirconia graded femoral ball-heads were developed. The potentialof this system follows from the properties of alumina (low wear, high hardness)and zirconia (high strength, high toughness). By combining alumina andzirconia, a functional gradient in hardness (high at the alumina surface) andtoughness (high in the zirconia-rich core) can be established. Due to thedifferent thermal expansion coefficients of Al2O3 and ZrO2, residual stressesare developed during cooling from the sintering temperature, which stronglyinfluence the mechanical properties like strength and toughness. Compressivesurface stresses in the outer alumina layer will also have a beneficial effecton wear resistance (Fig. 22.11) and strength (Fig. 22.12). Additionally, increasedtoughness can be observed in this kind of graded ceramic composite (Tilbrooket al., 2005).

22.5.3 WC/Co FGM materials

Another interesting group of structural gradient materials is that of cementedcarbides, commonly used, for example, for cutting tools and mining equipmentwhere high wear resistance and toughness are required. Using graded WC/

22.10 Overview of Al2O3 /TZP FGM ball-heads for hip prostheses,made by electrophoretic deposition and a schematic cross-section ofan FGM ball-head (Anné et al., 2004).


Co hardmetals, it is possible to improve both properties at the same time(Cherradi et al., 1994) (Fig. 22.13). This can be achieved by a gradation inthe binder content from the centre to the surface of the tool. The improvedmechanical properties of the FGM-cemented carbide are due to compressivestresses near the surface, which enable a reduction of the Co-phase contentfrom 6 to 3 wt% without any loss in apparent fracture toughness.

Two existing industrial technologies have been reported to produce gradedWC-Co hardmetals, i.e., liquid phase sintering of homogeneous WC-Co in acontrolled atmosphere or pressure assisted sintering of graded WC-Cocompacts. The first method is mainly use for nitride-containing WC-Co or

FGM disksAl2O3 disks

m = 9

288 MPa

m = 12

513 MPa

5.0 5.5 6.0 6.5In (strength)

2

1

0

–1

–2

–3

–4

In I

n (

1/S

)

22.11 Weibull plot of the biaxial strength (ISO 6474) of homogeneousAl2O3 and Al2O3/ZrO2 FGM disks with 80 vol% Al2O3 in the core, pureAl2O3 on the outside and intermediated convex graded profile. TheFGM disks were made by electrophoretic deposition and pressurelesssintering.

Ball wearFlat wear

10 175 241 260Residual compressive stress (MPa)

0.02

0.0180.0160.0140.012

0.010.0080.0060.0040.002

0

Wea

r vo

lum

e

22.12 Wear volume of step-graded Al2O3/ZrO2 FGM disks with 80vol% Al2O3 in the core and pure Al2O3 on the outside (Novak, et al.2005). The step-graded disks were made by sequential slip casting(Novak, et al. 2005).


WC-Ni cermets (Chen et al., 2000a; Narasimhan et al., 1995). The gradedCo distribution is formed during liquid-phase sintering through theestablishment of a nitrogen partial pressure gradient (Zackrisson et al., 2000;Chen et al., 2000b) or by a high-temperature carburization treatment of aWC–Co composite with sub-stoichiometric carbon content (Fischer et al.,1988). Pressure-assisted sintering techniques such as hot pressing, hot isostaticpressing and SPS can also be used to consolidate graded WC–Co compactsby solid state sintering (Tokita, 2003). Microwave sintering methods of WC/Co cemented carbides have also been reported (Willert-Porada et al., 1995).

Other graded cermet systems reported include Cr3C2/Ni (Seefeld et al., 1999),TiC/Ti (Cline, 1995), and TiC-Ni (Sabatello et al., 2000). The design conceptof the graded cermets is similar to that of WC/Co cemented carbide FGMs.

22.6 Future trends

Already a lot of scientific work has been established in processing gradedbulk ceramics and in modelling and optimization of the properties of thesematerials. Despite these technological successes, the number of commercialsuccesses is still limited because the production costs in most cases are stilltoo high compared to homogeneous materials. However, new productionprocesses are emerging on the market that allow economic production ofFGM bulk ceramics.

Many new target applications have already appeared or are rapidlyemerging,’ including the following:

∑ Thermo-mechanical applications (Kawai and Wakamatsu, 1995; Araki etal., 1994; Wakamatsu et al., 1999; Hofinger et al., 1999; Lee et al., 1996;Jiang et al., 1998)

22.13 Hardness–toughness relationship for WC-Co hardmetals(Cherradi et al., 1994).

FGM

Conventional hard metals

700 900 1100 1300 1500 1700 1900HV10

KlC

25

20

15

10

5


∑ Energy conversion (Niino, 1998; Mahan et al., 1997; Schilz et al., 1999;Sugiyama et al., 1998)

∑ Applications in optics and electronics (Wang et al., 1998; Palais, 1980;Komatsu et al., 1998; Le Goues et al., 1992, Miyamoto et al., 2005)

∑ Electrical and magnetic applications (Yamane et al., 1998; Nishida et al.,1999; Wu et al., 1996)

∑ Graded cemented carbide coatings on steel substrates (Put et al., 2003(c))∑ Biomedical applications (Kurzweg et al., 1998, Ban et al., 1999; Rogier

and Pernot, 1991, Anné et al., 2004).

22.7 Further reading

Since 1990, an international congress on Functionally Graded Materials isorganized by the International Advisory Committee of FGM (IACFGM)every two years. The proceedings of these congresses give a lot of informationabout research on FGM materials. The symposia on FGMs were held inSendai (1990), San Francisco (1992), Lausanne (1994), Tsukuba (1996),Dresden (1998), Estes Park, Colorado (2000), Beijing (2002) and Leuven(2004). The next one will be held in Chicago in 2006.

Due to their comprehensive description of the design, modelling, processing,and evaluation of FGMs as well as the many applications described, thefollowing books are recommended for further reading:

∑ Fundamentals of Functionally Graded Materials, by S. Suresh and A.Mortensen, A, Institute of Materials, Published by Woodhead, Cambridge,UK, 1998.

∑ Functionally Graded Materials: Design, Processing and Applications,edited by Y. Miyamoto, W.A. Kaysser, B.H. Rabin, A. Kawasaki and R.G.Ford, published by Kluwer Academic Publishers, Boston/Dordrecht/London,1999.

Several review articles have been written on FGMs, covering applications(Cherradi et al., 1994), processing (Mortensen and Suresh, 1995; Neubrandand Rödel, 1997), modelling (Markworth et al., 1995; Li et al., 2000) andfracture mechanics (Erdogan, 1995; Tilbrook et al., 2005).

In Japan, the FGM database http://fgmdb.nal.go.jp, http://fgmdb.nal.go.jp/e_whatsfgm.html is established. This database gives a good overview of theFGM concept, some examples and an overview of FGM literature. Additionallyan overview of almost all research groups involved in the FGM concept canbe found.

22.8 References

Anné, G., Vanmeensel, K., Vleugels, J., (2004), ‘Electrophoretic deposition as a near net


shaping technique for functionally graded biomaterials’, Mat. Sci. Forum, 492–493,213–218.

Araki, M., Sasaki, M., Kim, S., Suzuki, S., Nakamura, K., Akiba, M., (1994), ‘Thermalresponse experiments of SiC/ and TiC/C functionally gradient materials as plasmafacing materials for fusion applications’, J. Nucl. Mat., 212–215, 1329–1334.

Ban, S., Hasegawa, J., Maruno, S., (1999), ‘Fabrication and properties of functionallygraded bioactive composites comprising hydroxyapatite containing glass coated titanium’,Mat. Sci. Forum., 308–311, 350–355.

Bao, G., Cai, H., (1997) ‘Delamination cracking in functionally graded coating/metalsubstrate system’, Acta Mater, 45(3) 1055–1066.

Becker, T.L., Cannon, R.M., Ritchie, R.O., (2000), ‘An approximate method for residualstress calculation in functionally graded materials’, Mech. Mat., 85–97.

Bernhardt, R., Meyer-Olbersleben, F., Kieback, B., (1999), ‘The influence of hydrodynamiceffects on the adjustment of gradient patterns through gravity sedimentation ofpolydisperse particle systems in newtonian and viscoelastic fluid’, Mat. Sci. Forum,308–311: 31–35.

Cai, P.Z., Green, D.J., Messing, G.L., (1997a), ‘Constrained densification of alumina/zirconia hybrid laminates. 1. Experimental observations of processing defects’, J. Am.Ceram Soc., 80(8), 1929–1939.

Cai, P.Z., Green, D.J., Messing, G.L., (1997b) Constrained densification of alumina/zirconia hybrid laminates. 2. Viscoelastic stress computation’, J. Am. Ceram. Soc.,80(8), 1940–1948.

Chen, C., Awaji, H., (2003), ‘Transient and residual stresses in a hollow cylinder offunctionally graded materials’, Mat. Sci. Forum., 423–425, 665–670.

Chen, L.M., Lengauer, W., Ettmayer, P., Dreyer, K., Daub, H.W., Kassel, D., (2000a),‘Fundamentals of liquid phase sintering for modern cermets and functionally gradedcemented carbonitrides (FGCC)’, Int. J. Refract. Met. Hard Mat., 18(6), 307–322.

Chen, L.M., Lengauer, W., Dreyer, K., (2000b), ‘Advances in modern nitrogen-containinghard-metals and cermets’, Int. J. Refract. Met. Hard Mat., 18(2–3), 153–161.

Cherradi, N., Kawasaki, A., Gasik, M., (1994), ‘Worldwide trends in functional gradientmaterials research and development’, Compos. Eng., 4(8), 883–894.

Chu, J., Ishibashi, H., Hayashi, K., Takebe, H., Morinaga, K., (1993), ‘Slip casting ofcontinuous functionally gradient material’, J. Ceram. Soc. Japan, 101, 818–820.

Cichocki, F.R., Jr, Trumble, K.P., (1998), ‘Tailored porosity gradients via colloidal infiltrationof compression-molded sponges’, J. Am. Ceram. Soc., 81(6), 1661–1664.

Cline, C.F., (1995), ‘Preparation of gradient TiC cermet cutting tools’, in Ilschner, B,Cherradi, N., (eds), Proc. 3rd Int. Symp. on Structural and Functional Gradient Materials,Presses Polytechniques et Universitaires Romandes, Lausanne, 595.

Erdogan, F., (1995), ‘Fracture mechanics of functionally graded materials’, Mater. Res.Soc. Bull, 20(1), 43–44.

Finot, M., Sureshs (n.d.), http://ninas.mit.edu/lexcom/www/multitherm.html.Fischer, U.K.R., Hartzell, E.T., Akerman J.G.H., (1988), ‘Cemented carbide body used

preferably for rock drilling and mineral cutting’, US Patent No. 4,743,515, 10 May1988.

Fukui, Y., Takashima, K., Ponton, C.B., (1994), ‘Measurement of Young’s modulus andinternal friction of an in situ Al–Al3Ni functionally gradient material’, J. Mat. Sci., 29,2281–2286.

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597

Index

absorption frequency 470–1accommodation processes 434, 439–45

GBS with diffusional flow 439, 440GBS with dislocation movement 439,

440–2shear thickening creep 444–5solution-precipitation creep 442–4

activation energy 264–5, 437–8active oxidation 268additives

fugitive 515–18magnesia-spinel refractories 366–7sintering see sintering additives

advanced carbon materials, nanostructuredcoatings on see nanostructuredcoatings

advanced ceramics: basic material properties335, 336

aerospace industry 93–4alkali attack 360–1, 367, 368, 391–2Almax alumina fibre 64, 66a-alumina 62a-b phase transformation 39a:b ratio 161–6, 167a/b-sialon 155, 157, 160, 494–8, 510

aspect ratio effect 497composition design 495–6thermal and Young’s modulus mismatches

497–8and transformation 493volume fraction of elongated grains

496–7a-sialon 155–6, 157, 158–9, 492–3, 522,

537–8, 556–7change in the amount in tape laminated

sialon FGMs 168eroded surfaces 544–5erosion performance 543–4microstructure and erosion mechanisms 553microstructure and mechanical properties

550–2

powder bed methods for sialon FGMs170–2

a-silicon nitride 155, 156, 475carbon-ceramic alloys 522–3, 526, 527

alumina 3–4, 536duplex microstructures with YAG 118fibres suitable for reinforcement of glass/

glass-ceramic matrix composites62–6

material properties 336particulate composites 107–9sintering additive 34, 40–1, 44, 50–1, 53sintering in magnesia-alumina system

364thermal shock behaviour 413, 414, 415

reinforced with iron particles 414, 415resistance parameters 411

alumina-aluminium titanate composite114–15

alumina-based FGMs 587–8alumina-CNT composites

mechanical properties 320–3, 324–5nanoceramics 251–3

alumina-lanthanum phosphate (monazite)335–44, 352

machinability 341–4, 345mechanical properties 338–41microstructure 341, 342, 343

alumina-matrix LGMs 138–50alumina/aluminium titanate and alumina-

zirconia/aluminium titanate systems140–7

alumina/mullite and mullite/ZTA/mullitesystems 138–40

alumina/mullite/AT hybrid 147–50liquid infiltration processing 138

alumina membrane 324–5alumina-molybdenum nanocomposites 246alumina-nickel nanocomposites 246alumina-rich magnesia-spinel 360alumina and silica mix fibre 63, 64

Index598

alumina-silicon carbide nanocomposites101–2, 103, 114–16, 118–19,122–4, 125, 244–5

alumina/Ti3SiC2 system 351alumina-titanium carbide cutting tools 121,

124alumina/zirconia FGM disks 580–1, 582, 587,

588, 589alumina/zirconia nanoceramics 253–5aluminate platelets 4aluminium nitride

AlN-BN powder beds 171–3AlN polytypoids 537–8AlN/TiN multilayers 234–6crystal forms 235–6dewetting behaviour 468

aluminium titanate 3–4, 114–15, 125alumina-matrix LGMs

alumina/aluminium titanate andalumina-zirconia/aluminium titanatesystems 140–7

alumina/mullite/AT hybrid 147–50thermal shock resistance parameters 411

aluminium/titanium boride multilayer 228, 229aluminosilicate glass 89–92analytical models for thermal stresses 584–7anisotropy 427

and Hamaker constants 471–2anorthite 81apatites 564, 567, 569apparent fracture toughness 181–5, 210–11

calculations of 185–7silicon nitride laminates 194–9

aragonite bricks 1arc-discharge evaporation 311archery bows 2armchair-type SWCNTs 310Ashby-Verrall (A-V) model 440aspect ratio 51–2

a/b-sialon 497critical 89, 90and fracture toughness of sialons 551

atomic force microscopy 483atomistic simulations 474–5attritor milling 524–6

B-phase Y-Si-Al-O-N system 564–5, 567ball milling 519, 524–6ballistic tests 203–4, 205, 208–9bamboo 1, 2barium-osumilite matrix composite 87–9barium-stuffed cordierite matrix composite

87–9barium titanate 314beryllium oxide 411b-grain seeds 39–40b-sialon 155–6, 157, 158, 492

powder bed method for sialon FGMs 170–3b´-sialon layer 40–1b-silicon nitride 155, 156

carbon-ceramic alloys 522–3, 526, 527crystallography of interphase boundaries

475–7, 479–82elongated grains 538–9

biaxial stress state 585–7bifurcation, crack 187–9, 200, 201, 202binding energy 446, 453bioglass 518–19biomedical implants 94, 518–19, 587–8Biot modulus 401, 406–7bone 1–2bone-substituting implants 518–19borides 15boron carbide 71boron carbide-based laminates 201–9

ballistic testing 203–4, 205, 208–9boron carbide/boron carbide–SiC laminate

for increased fracture toughness210–11

design and mechanical behaviour 201–3microstructures of three-layered boron

carbide/boron carbide-SiClaminates 206–9

processing of boron carbide-SiC laminates189–93

undesirable influence of tensile residualstresses 204–6

boron nitride 71cell boundary 15cubic 216h-BN see h-BNreinforced nanocomposites 246–8

borosilicate glasses 89–92bridging 74Burgers’ vectors 220–1

calcined spinel 363calcium, as dopant 474calcium a-sialon 537–8, 54–5, 550–2, 553,

556–7calcium aluminosilicate (CAS) 418–19, 420–1,

423–4calcium oxide 366–7carbon 260

material properties 336nanostructured coatings on advanced

carbon materials see nanostructuredcoatings

carbon black 515–18hot pressing of carbon-ceramic alloys

528–30carbon–ceramic alloys 514–35

carbon as fugitive additive 515–18comparison of silicon nitrides with carbon

Index 599

additions prepared by HIP andpressureless sintering 518–23

in-situ processing of silicon nitride-siliconcarbide composites by carbonaddition 524–30

silicon nitride ceramics reinforced withcarbon fibres and carbon nanotubes530–3

carbon fibre-reinforced glass/glass-ceramicmatrix composites 95

carbon fibre-reinforced sialons 500–5effects of processing parameters 500–2thermal mismatch 502–4

compensation in thermal mismatch504–5

carbon fibre-reinforced silicon nitride ceramics530–3

carbon nanotubes (CNTs) 515alumina-CNT composites see alumina-

CNT compositessilicon carbide coating 260, 265–70, 281

growth mechanism 266–70microstructure 265–6oxidation resistance 271–3silicon carbide–coated CNTs/silicon

carbide composites 277–80silicon carbide-coated CNTs/WC-Co

composites 275–7, 281silicon nitride ceramics reinforced with

530–3structure, synthesis and properties 309–12

carbon nanotubes-ceramic composites 309–33future trends 329–30preparation 313–19

densification 317–19in situ synthesis of the ceramic in CNT

samples 314–15in situ synthesis of CNTs in ceramic

powder 315–17mixing CNTs with ceramic powders 313

properties 320–9electrical 325–8, 330mechanical 320–5, 330thermal 328–9, 330

catalytic chemical vapour deposition (CCVD)311, 312, 315–16

cell boundaries 11–12, 14–15, 28–9cells 11–12, 14–15, 28–9cement kilns 359–62, 376, 390–2

rotary see rotary cement kilnscemented carbide (WC-Co)

functionally graded materials 588–90silicon carbide-coated carbon nanotubes/

WC-Co composites 275–7, 281silicon carbide-coated diamond/WC-Co

composites 273–5, 281centrifugal casting 578, 579

ceramic matrix materials 80Ceran 81cermet nanocomposites 285–308

future trends 304microstructure 290–300

evolution and control 292–4interfaces 290, 294–6, 301–2particle occlusion 296–8residual stresses and fracture strength

299–300processing 285–90properties 300–4

functional 302mechanical 300–2oxidation resistance 302–4

cermets 80, 285see also cermet nanocomposites

cetyltrimethylammonium bromide 314channel cracking 201, 202charge effect 447–8, 453chemical attack 360–1, 367, 368, 391–2chemical vapour deposition (CVD) 68–70chemical vapour infiltration (CVI) 421–3chevron-notched specimen flexural technique

(CN-technique) 409chromium carbide 108–9, 125chromium ions 360–1cleavage steps 207, 208–9coating loss 368Coble diffusional creep 436, 438coefficient of thermal expansion (CTE) 178–9,

180–1, 183boron carbide laminates 201, 203fibre-reinforced glass/glass ceramic matrix

composites 82–3oxynitride glasses 566silicon nitride laminates 193thermal shock and fibre-reinforced CMCs

425whisker-reinforced silicon nitride ceramics

42–3coextrusion 11–12coherency stresses 219–22coking 72cold pressing 73–4columnar grain boundaries 232–6combustion synthesis 583co-milling 313compositional FGMs 577–8compound machinable ceramics 335compression: and tension in creep 437–9compression axis 450–1, 452compressive stresses, residual 194–9computer modelling of interfaces 483conchoidal fractures 547–8, 555–6conductivity, electrical 325–8conductivity, thermal 43, 328–9

Index600

consolidation processes 578–81continuous fibre reinforcements 60–1copper/titanium nitride multilayer 228, 229cordierite (cristobalite) 81Corningware 81corrosion 371–2crack bifurcation 187–9, 200, 201, 202crack blunting 10crack bridging 147, 536

particulate composites 104, 107–9, 110, 112silicon nitride ceramics 555

whisker-reinforced 38, 41–2crack deflection 10, 19, 536

laminate toughening and residual stresses187–9

machinability of alumina-lanthanumphosphate 343–4, 345

particulate composites 104, 105–7whisker-reinforced silicon nitride ceramics

41, 49–50crack delamination 19–20, 21, 23, 28crack initiation

magnesia-spinel ceramic composites 373,384–5

particulate composites 104–5thermal shock 410–11

crack kinking 187–9crack length 181–2, 187

and apparent fracture toughness 195–6crack pinning 352, 353crack propagation

FMs 19–20thermal shock resistance 26–8

magnesia-spinel ceramic composites 374,385–6, 389–90

and residual stresses in silicon nitride-based laminates 195–9

stable and unstable crack growth 197–9thermal shock 410–12whisker-reinforced silicon nitride ceramics

46–7cracking/cracks

channel cracking 201, 202FGMs and residual stresses 581–2FMs and thermal shock resistance 26–8intergranular cracks 382–3, 384, 545, 553,

557thermal shock and matrix cracking in fibre-

reinforced CMCs 417–19, 422–4transgranular cracks see transgranular

crackssee also microcracking

creepaccommodation processes controlling

superplasticity 439–45creep resistance of whisker-reinforced

silicon nitride ceramics 53–4

liquid phase-enhanced 250–1oxynitride glasses 570

creep rate 564parameters and superplasticity 437–8particulate composites 118–19

cristobalite (tridymite) 81critical aspect ratio 89, 90critical dimension 407critical flaw size 111, 301, 381–2, 389critical particle size 287–8critical quench temperature 372critical quenching temperature difference

407–8, 411–12, 416–17, 425–7critical tensile stress 206critical thermal shock-induced stress 426cross-ply CMCs 420–1crystallinity

crystal structures of spinels 362crystalline phase 38–9crystalline size of silicon carbide

nanostructured coatings272–3

oxynitride glasses 563–6, 567, 568–70crystallography of interphase boundaries

475–82cubic structure 235–6cutting tools 55, 94, 121, 124

damage zones 143–7debond cracks, thermal 418–19, 420, 423–4Debye length 447–8deflocculants 286deformation processes 232–6degree of damage 374, 387–8delamination cracks 19–20, 21, 23, 28densification

CNT-ceramic composites 317–19FGM powder compacts 581–3magnesia-spinel ceramic composites

363–5see also under individual methods

densityoxynitride glasses 568silicon nitride/h-BN composites 346silicon nitride and silicon carbide-based

ceramics 540, 541depth profiles 142–3, 148–9de-sintering 517–18dewetting behaviour 467–9diamond 216

silicon carbide coating nanostructured 260,261–5, 281

oxidation resistance 270–3silicon carbide-coated diamond/WC-Co

composites 273–5, 281diffuse dark field imaging 463, 464–6diffusion

Index 601

between layers of sialon FGMs 164–5cermet nanocomposites 298GBS accompanied by 439, 440grain boundary diffusion 249–50, 436, 440lattice diffusion 436, 440superplasticity and 434

improvement of processes controllingdiffusion 446–8

diffusional creep 249, 436, 438dihedral angle 293–4dimensionless stress/dimensionless time plot

404, 405dip-coating 12, 13–14discontinuous fibre reinforcements 60–2disilicates 564, 567dislocations 219

GBS accompanied by dislocationmovement 439, 440–2

GBS and superplasticity 453hardening

changes in dislocation line energy 219,222–6

coherency stresses 219–22lateral flow of material 219, 227–30

machinable nanocomposite ceramics 352dislodgement of grains 548, 555–6dispersants 101dolomite-based refractories 361, 367–8, 391dopants 286double-walled CNTs (DWCNTs) 311, 312

see also carbon nanotubes (CNTs)drag force, for a particle 293–4, 297–8drilling rate 343, 344dry processing 578duplex microstructures 3–4, 118dynamic model for erosion 552dynamic Ostwald ripening process 514

edge-defined film-fed growth (EFG) 64, 67effective pore diameter 135–6elastic mismatch 222–6, 497–8elastic modulus see Young’s moduluselastic strains 402–3electric field, and diffusion 447–8electrical properties

CNT-ceramic composites 325–8, 330CNTs 312

electrofused spinel 363electron density 453electron energy-loss spectroscopy (EELS) 467electron microscopy 482–3electrophoretic deposition (EPD) 578, 579–81elongated grains see grain elongationenergy absorption

fibrous monolithic ceramics 19–20and plastic deformation 58

energy dispersive X-ray (EDX) 463

energy-loss near-edge structures (ELNES)467

engineering ceramics 99enstatite 81equi-biaxial stress states 585–7equilibrium amorphous films 294–6equilibrium film thickness 295, 469–72

effect of film composition on 473–5eroded surfaces 544–50erosion see wearerosion rate, steady 543–4, 552expansion mismatch see thermal expansion

mismatch

Faber’s theory 49failure of fibrous monolithic ceramics

and high temperature 21–2, 23mechanisms 16–18

fast consolidation techniques 256fast cooling method 173–4fast heating tests 407–8fatigue

behaviour of glass-ceramic matrixcomposites 85–7

resistance of whisker-reinforced siliconnitride ceramics 54

feedrod 11, 12–13ferroelasticity 254ferroelectric ceramic matrices 328ferroelectric-domain switching 253–4fibre-reinforced ceramic matrix composites

4–5, 10, 178thermal shock 416–8

thermal shock damage and mechanicaland thermal properties 417–25

theoretical considerations 425–7fibre-reinforced glass/glass-ceramic matrix

composites 58–98application areas 93–5future trends 95methods for manufacture 72–80, 82–3microstructural observation 92–3properties 80–92

fibre-reinforced glass/glass ceramicmatrix composites 81–5

glass-ceramic matrix composites 85–7silicon carbide whisker-reinforced glass

and glass-ceramic composites87–92

types of fibre suitable as reinforcements60–72

non-oxide fibres 60–1, 68–72oxide fibres 60–1, 62–7

fibre-reinforced sialons 510carbon fibre-reinforced 500–5silicon carbide fibre-reinforced 505–9

fibrils, mineralised 1–2

Index602

fibrous monolithic ceramics (FMs) 9–32future trends 28–9history 9–11mechanical properties 15–28

high-temperature properties 21–8room-temperature properties 16–20

processing 11–14structures 14–15

architectures 15, 16material combinations 14–15

fibrous preforms 77–8, 79figures of merit for wear 120–1filament winding 77–8, 79film thickness

equilibrium film thickness see equilibriumfilm thickness

at interfaces 463–7finite element analysis 585flaw tolerance 3–4flexural strength

alumina-lanthanum phosphate 338–9alumina-matrix LGMs 139carbon-ceramic alloys 518, 523, 529, 531fibre-reinforced glass/glass-ceramic matrix

composites 89, 90fibrous monolithic ceramics 16–18

high temperature 21, 22–6oxynitride glasses 566silicon nitride/h-BN 346–8silicon nitride and silicon carbide-based

ceramics 542whisker-reinforced silicon nitride ceramics

43–6flexural tests 409flow stress 218–19FP fibre 62fracture mechanics 411–12fracture steps 201, 202fracture strength

cermet nanocomposites 299–300, 301CNT-ceramic composites 320, 323oxynitride glasses 570–1thermal shock and 423

fracture stress test 46fracture surface energy 375–6, 381, 389fracture surfaces

alumina-lanthanum phosphate 341, 343boron carbide laminates 207, 208carbon-ceramic alloys 519–22carbon fibre-reinforced sialons 501silicon nitride/h-BN 350–1silicon nitride laminates 199–201, 202see also microstructure

fracture toughness 412alumina-matrix LGMs 139–40, 149–50apparent see apparent fracture toughnesscemented carbide FGMs 588–90

CNT-ceramic composites 320–4, 330fibre-reinforced glass/glass ceramic matrix

composites 87–9increasing by residual stresses control see

residual stresses controlmagnesia-spinel ceramic composites 375,

381, 389nanostructured silicon carbide coatings

274–5, 278–9oxynitride glasses 569sialons 540–1, 542, 551–2

carbon fibre-reinforced sialoncomposites 504, 505

silicon nitride and silicon carbide ceramics536, 540–1, 542

freeze-gelation method 95Fresnel defocus imaging 463, 464–7fringe spacing, Fresnel 465, 466fugitive additives 515–18functional particulate composites 124functional properties 302functionally graded materials (FGMs) 154–5,

451, 575–96applications 577, 590–1classification 577concept 575–6design for structural applications 584–90

alumina and zirconia-based FGMs587–8

thermal stress calculations 584–7WC/Co FGMs 588–90

functionally graded sialon ceramics160–75

production techniques 161–74future trends 590–1processing 577–83

densification of FGM powder compacts581–3

shaping and consolidation processes578–81

sialon-based see sialon-based functionallygraded materials

see also layered-graded materials (LGMs)fused deposition of ceramics (FDC) 29

gas pressure sintering (GPS) 158gas turbines 54–5, 400Gaussian focus 466gel–casting 289glass 81

fibre–reinforced glass/glass-ceramic matrixcomposites see fibre-reinforcedglass/glass-ceramic matrixcomposites

grain boundary glass 551–2, 553, 556glass-ceramic composites 81

oxynitride glasses 560–74

Index 603

crystallisation to form multi-phaseglass-ceramic composites 562–9

properties 85–7glass doping 296–7glassy materials 323, 324glassy phases 441–5, 447graded cermet systems 590gradient: classification of FGMs by 577grain boundaries 290, 461

role in superplasticity 453shear along in thin films 232–6

grain boundary diffusion 249–50, 436, 440grain boundary glass 551–2, 553, 556grain boundary microcracking 371grain boundary pinning 293grain boundary sliding (GBS) 117–18, 249,

250superplasticity 434, 436–7, 438, 449

electron density and binding energy 453GBS with diffusion 439, 440GBS with dislocation movement 439,

440–2types of defects created 453

grain debonding 144grain dislodgement 548, 555–6grain ejection 544, 553, 557grain elongation

aspect ratio in a/b-sialon 497and erosion resistance 556silicon nitride 538–9volume fraction in a/b-sialon 496–7

grain growth 293–4, 365grain refinement 148, 292–3, 300–1grain size 44, 45, 541

magnesia and strength of magnesia-spinelceramic composites 378–9

and superplasticity 437–8reduction 445–6, 447–8

grain size FGMs 577graphite, in carbon-ceramic alloys 515–18,

528–30gravitational sedimentation 578, 579growth mechanisms

silicon carbide coating on carbonnanotubes 266–70

silicon carbide coating on diamondparticles 262–5

h-BN 246–8material properties 335, 336, 337silicon nitride/h-BN machinable

nanocomposites 344–51Hall-Petch relation 248Hamaker constant 470–2, 473–4, 482hard nanoparticle dispersed nanocomposites

244–5hardness

alumina-lanthanum phosphate 338, 339–41alumina–matrix LGMs 139–40, 143,

149–50carbon-ceramic alloys 523, 528, 529CNT-ceramic composites 322, 323multilayered ceramics see multilayered

ceramicsnanoceramics 248nanostructured silicon carbide coatings

274–5, 276–9oxynitride glasses 566sialon FGMs 163, 164, 165, 166, 168, 169sialons 540, 542, 551silicon nitride/h-BN 346–8silicon nitride and silicon carbide-based

ceramics 540, 542WC-Co FGMs 588–90

hardness indentation test 46, 143–7heat transfer coefficient (HTC) 401, 408heat treatment 304

silicon carbide fibres for silicon carbidefibre-reinforced sialon 506–7

helical vector 310, 311helicity 310, 312Hertzian indentation 144–7heterophase boundaries 461high-energy ball-milling (HEBM) 255high-performance ceramics 58high-pressure sintering 255–6high resolution transmission electron

microscopy (HRTEM) 463, 464–7high strain rate superplasticity (HSRS) 454high-temperature gas turbines 54–5hip replacement prostheses 588homophase boundaries 461hot isostatic pressing (HIP) 36

CNT-ceramic composites 319compared with pressureless sintering for

carbon-ceramic alloys 518–23sialons 157–8

hot pressing (HP) 36, 37, 190, 191–2CNT-ceramic composites 317–19fibre-reinforced glass/glass ceramic

composites 74–5, 76micro- and nano-carbon added carbon/

silicon nitride composites 524–30sialons 157–8thermal shock behaviour of hot-pressed

alumina 414, 415reinforcement with iron particles 414,

415hot strength 371, 376, 377HPZ (silicon carbonitride) 71hybrid CMCs 4–5hybrid extrusion 13–14

impregnation by a precursor 316–17

Index604

impurities, and superplasticity 446, 453in-situ processing 3

silicon nitride/silicon carbide compositesby carbon addition 524–30

spinel based refractories 365–6synthesis of CNT-ceramics

ceramic in CNT samples 314–15, 329CNTs in ceramic powder 315–17, 329

whisker-reinforced silicon nitride ceramics33

indentation hardness tests 46, 143–7indentation prints 279indentation thermal shock tests 409indium gallium arsenide 221–2induced stress field 401–7, 427infiltration process 131, 150–1

infiltration kinetics and characteristics132–7

modelling infiltration kinetics 133–4studies on infiltration kinetics 134–7

processing of LGMs 137–8infrared spectroscopy 515, 516interfaces

cermet nanocomposites 290, 294–6, 301–2degradation and thermal shock 422–3,

424–5non-oxide ceramic composites 461–90

accuracy of TEM for measurement offilm thickness 463–7

crystallography of interphaseboundaries 475–82

effect of intergranular film compositionon equilibrium film thickness473–5

equilibrium film thickness 469–72future trends 482–3wetting behaviour of interphase

boundaries 467–9sharpness 226

interfacial diffusion 298interfacial fracture resistance 19–20intergranular cracks 382–3, 384, 545, 553, 557intergranular films 294–6, 304, 461–3

accuracy of TEM techniques formeasurement of thickness 463–7

effect of film composition on equilibriumfilm thickness 473–5

equilibrium film thickness 469–72interlocking effects 553, 554, 556intermetallic compounds 304internal friction 423interphase boundaries 461

crystallography of 475–82equilibrium film thickness 469–72wetting behaviour 467–9

inverse spinels 362inverted Hall-Petch model 248

inviscid melt technique 67iron aluminate 315iron nitride 314–15iron oxide 366–7island structure 442–3isostatic pressing 73

hot see hot isostatic pressing (HIP)isostrain 450–1, 452isostress 450–1, 452Iw-phase Y-Si-Al-O-N system 564–5

Jänecke prisms 156, 561–2, 563joining ceramics 449–52‘joining’ temperature 180, 183joint prostheses 588

kinking, crack 187–9Kozeny-Carman equation 133–4

laminated composites 10, 11, 510ceramic-metal systems 401design for toughness by residual stresses

control see residual stresses control‘LAMINATES’ project 210lamination techniques for sialon FGMs 161–70

powder lamination 161–6tape lamination 166–70

lanthanum phosphate (monazite)alumina-lanthanum phosphate 335–44, 352material properties 335, 336, 337

Lanxide process 75, 77laser-heated floating zone method 66, 67laser sublimation of graphite rods 311lateral flow of material, hardening due to 219,

227–30lattice diffusion 436, 440lattice rotations 232, 233layer thickness 182–3

thickness ratio of layers 183–5layered ceramics 4, 401layered-graded materials (LGMs) 131–53

characterisation and properties of alumina-matrix LGMs 138–50

alumina/aluminium titanate andalumina-zirconia/aluminium titanatesystems 140–7

alumina/mullite and mullite/ZTA/mullite systems 138–40

alumina/mullite/AT hybrid 147–50infiltration kinetics and characteristics

132–7infiltration processing 137–8

layered or soft materials 335, 336lead zirconate titanate (LZT) 80Lifshitz grain boundary sliding 436liquid infiltration see infiltration processliquid phase 34–5

Index 605

liquid phase-enhanced creep 250–1liquid phase sintering

magnesia-spinel system 365WC-Co FGMs 589–90

lithium aluminosilicate (LAS) 81, 417–18glass-ceramic matrix composites 85–7sintering additive and fibre-reinforced

sialons 500, 507–8, 509Ln-Si-Al-O-N system 562, 567–8Ln-Si-O-N system 562, 567–8load-displacement curve 501–2

machinable nanocomposite ceramics 334–55alumina–lanthanum phosphate 335–44

experimental procedure 338machinability 341–4mechanical properties 338–41microstructure 341

design principles of machinable ceramics334–5, 336, 337

machinable nanocomposites 351–2, 353silicon nitride/h-BN 344–51

experimental procedure 344–5mechanical properties 346–8microstructure and machinability

348–51machining flaws 571macro-scale engineering 29macrostresses 299–300magnesia

grain size and strength of magnesia–spinelceramic composites 378–9

sintering of 363–4sintering additive 34–5, 44, 53–4, 286sintering in magnesia-alumina system 364thermal shock resistance parameters 411

magnesia-chrome refractory materials 360–1,391

magnesia-rich spinel bricks 359, 360magnesia-spinel ceramic composites 359–99

crystal structures 362densification 363–5future trends 390–2in-situ formed/preformed spinel based

refractories 365–6industrial applications and properties

366–72performance 370–2rotary cement kilns 367–70, 391type and effect of additives 366–7

mechanical properties and thermal shockbehaviour 375–89

comparison of fraction retainedstrength and R´´´´ parameter 388–9

critical defect size 381–2, 389degree of damage 374, 387–8fracture surface energy 375–6, 381, 389

fracture toughness 375, 381, 389R and R´´´´ parameters 384–6relative strength 386–7strength and magnesia grain size 378–9work of fracture 376, 382–4, 389Young’s modulus 370, 376, 377,

379–80production of spinel 362–3thermal shock 359–60, 372–4, 391–2

magnesia-spinel clinker, sintered 363, 366magnesium aluminate 315–16magnesium aluminate spinel see magnesia-

spinel ceramic compositesmagnetoresistance properties 302mass gain 264mass loss 343, 344matrix cracking stress 426maximum tensile stress 404–6mechanical properties 3–4

advanced ceramics and layered or softmaterials 335, 336

alumina-lanthanum phosphate 338–41alumina-matrix LGMs 139–40, 141,

146–7, 149–50boron carbide-based laminates 201–3carbon-ceramic alloys 523, 528–30, 531, 533cermet nanocomposites 300–2CNT-ceramic composites 320–5, 330CNTs 312fibre-reinforced glass/glass-ceramic matrix

composites 80–92fibrous monolithic ceramics 15–28magnesia-spinel composite refractory

materials 370–1, 375–89nanostructured coatings 274–5, 276–9oxynitride glasses 566, 567–8, 570–1silicon nitride-based laminates 193–4silicon nitride/h-BN machinable

nanocomposites 346–8silicon nitride and silicon carbide-based

ceramics 540–1, 542microstructure and 550–2

thermal shock damage and for fibre-reinforced CMCs 417–25

whisker-reinforced silicon nitride ceramics43–54

melilite 522, 567metal-matrix composites (MMCs) 3, 59, 60metal nanoparticle dispersed nanocomposites

245–6metal-reinforced ceramic matrix

nanocomposites see cermetnanocomposites

metals 58indentation damage 144

mica 336microcomposites, particulate

Index606

room-temperature strength 110–14wear 120–2

microcrack toughening 104, 109, 112microcracking

carbon fibre-reinforced sialon composites502

grain boundary microcracking 371magnesia-spinel ceramic composites

378–89spontaneous in particulate composites 105,

106microcracking density 423microfabrication by coextrusion (MFCX)

12–13microhardness 52

nanostructured coatings 276–9oxynitride glasses 568whisker-reinforced silicon nitride ceramics

52micro-nano type ceramic composites 243,

244–8, 257hard nanoparticle dispersed

nanocomposites 244–5metal nanoparticle dispersed

nanocomposites 246–8soft and weak nanoparticle dispersed

nanocomposites 246–8micro-scale engineering 28–9microstresses 299–300

thermal 103–5, 110microstructure 536–59

a/b-sialon composites 496, 497alumina-CNT nanoceramics 253alumina-lanthanum phosphate 341, 342,

343alumina-matrix LGMs 140–1boron carbide laminates 206–9carbon-ceramic alloys

CNT-reinforced 531–3HIP compared with pressureless

sintering 519–23hot pressing 524–8

cermet nanocomposites 290–300changes due to making a multilayer

230–6eroded surfaces 544–50

sialon 544–5silicon carbides 547–50silicon nitride 546–7

and erosion mechanisms 552–6sialon 553silicon carbides 555–6silicon nitride 553–5

erosion performance 543–4FGMs 580–1fibrous monolithic ceramics 14–15, 16material characterisation 540–3

and mechanical properties 550–2microstructural development of particulate

composites 100–3observing 92–3refinement and superplasticity 445–6sialons 537–8, 550–2silicon carbide nanostructured coatings

261–2, 265–6, 273–4, 279–80silicon carbides 539–40silicon nitride 538–9silicon nitride/h-BN machinable

nanocomposites 348–51silicon nitride laminates 199–201, 202tough interlocking microstructure of

elongated grains 514wear resistance of particulate composites

121–2whisker-reinforced silicon nitride ceramics

38–41microwave sintering 256, 583milling 101, 190, 191mineral crystals 1–2misfit dislocations 219–22mismatch

elastic 222–6, 497–8in shrinkage 101thermal expansion see thermal expansion

mismatchmobility of a particle/pore 297–8modulus ratio 60modulus of rupture (MOR) 87–9molybdenum/NbN multilayer 228, 229monazite see lanthanum phosphatemonolithic ceramics 124–5

accommodation processes controllingsuperplasticity 439–41

fibrous see fibrous monolithic ceramicshigh-performance 58thermal shock 400, 410–13

monolithic films 231–2monticellite 361mullite 81, 569

alumina-matrix LGMsalumina/mullite and mullite/ZTA/

mullite systems 138–40alumina/mullite/AT hybrid 147–50

multiaxial architecture 15, 16multifilament coextrusion 12–13multifunctional layer composites 151multilayered ceramics 216–40

behaviour of multilayer structures 217–19future trends 237grain size, superplasticity and 449–52hardening mechanisms 219–30

changes in dislocation line density 219,222–6

coherency stresses 219–22

Index 607

lateral flow of material 219, 227–30microstructural changes due to making a

multilayer 230–6deformation processes and

microstructure of the film 232–6hardening due to internal stresses

230–2Multitherm program 585multi-walled carbon nanotubes (MWCNTs)

252, 309–11, 312hot-pressing of MWCNT-ceramic

composites 318–19silicon carbide coating on 265–70, 281

silicon carbide-coated MWCNTs/silicon carbide composites 277–80

silicon carbide-coated MWCNTs/WC-Co composites 275–7

see also carbon nanotubes (CNTs)

Nabarro-Herring diffusional creep 436nacre 1–2nanoceramics 243–4, 248–56, 257

fabrication 255–6high toughness and toughening mechanism

251–5superplasticity

high temperature 250–1low temperature 248–50

variety of hardness 248nanocomposite processing 514–15nanocomposites

carbon-added silicon nitride composites524–30

carbon nanotubes-ceramic composites seecarbon nanotubes-ceramiccomposites

machinable see machinable nanocompositeceramics

metal-reinforced see cermetnanocomposites

particulate compositescreep reduction 118–19room-temperature strength 114–16wear 122–4

silicon carbide particle-reinforced sialoncomposites 499

see also nanophase ceramic compositesnanocrystalline ceramics 446nanoindentation 232, 233nanophase ceramic composites 243–59

future trends 257micro-nano type 243, 244–8, 257nano-nano type (nanoceramics) 243–4,

248–56, 257nanostructured coatings 260–84

applications 273–80silicon carbide-coated carbon

nanotubes/silicon carbidecomposites 277–80

silicon carbide-coated carbonnanotubes/WC-Co composites275–7, 281

silicon carbide-coated diamond/WC-Cocomposites 273–5, 281

coating method of nanostructured siliconcarbide 261–73

coating assembly 261coating on carbon nanotubes 265–70,

271–3coating on diamond particles 261–5,

270–1oxidation resistance 270–3

nanotechnology 29natural composites 1–2Nd-Mg-Si-O-N system 568–9Nextel fibres 63, 65Niaproof 64Nicalon fibres 71

fibre-reinforced composites 417–19, 420–4nickel ferrate 314nickel spinel 302–3niobium nitride 217, 220, 222, 223, 224, 225–6

Mo/NbN multilayer 228, 229nitrogen 473–4nitrogen-containing lithium aluminosilicate

glass (NLAS) 507–8, 509non-destructive techniques 409–10non-oxide fibres 60–1, 68–72non-oxide fibrous monolithic ceramics 15non-wetting behaviour 467–9normal spinels 362nuclear reactors 400nucleation 563, 564

O-sialon ceramics 155–6, 157, 159occluded metal particles 290, 296–8occluded pores 290octahedral sites 362optical microscopy 93optical properties 302organic additives 313orientation

relationships and interphase boundaries475–82

of whiskers see whisker orientationOrowan’s equation for dislocation movement

443–4Ostwald ripening process 514outer layers, residual stresses in 194–9oxidation

active 268resistance of cermet nanocomposites 302–4resistance of FMs at high temperature 22–

4, 25

Index608

resistance of silicon carbide nanostructuredcoatings 270–3

oxide reduction 287–8oxide fibres 60–1, 62–7oxide fibrous monolithic ceramics 15oxidised starting powder 515–18oxynitride glasses 560–74

crystallisation to form multi–phase glass-ceramic composites 562–9

oxynitride glass-silicon carbide composites569–71

potential applications 560–1structure and properties 561–2see also glass-ceramic composites

parallel electron energy loss spectroscopy(PEELS) 463

partial pressure of oxygen 302–3partial sintering 515particle coarsening 292particle occlusion 290, 296–8particle shape 292particle size 571particulate composites 4–5, 99–128, 178

future trends 124–5high-temperature strength 116–20powder processing and microstructural

development 100–3room-temperature strength 110–16

ceramic nanocomposites 114–16microcomposites 110–14

sialon composites 498–9thermal microstresses 103–5, 110thermal shock 413–16toughening 105–10wear 120–4

microcomposites 120–2nanocomposites 122–4

perhydropolysilazane polymer (PHS) 71physical properties

alumina-matrix LGMs 139–40, 141,146–7, 149–50

silicon nitride and silicon carbide-basedceramics 540, 541


see also mechanical propertiesplasma activated sintering (PAS) 256plastic deformation 58plasticiser 190, 191Poisson’s ratio 180–1, 183, 203polycarbosilane 68, 69polyethyleneamine (PEI) 313polymer matrix composites (PMCs) 3, 59, 60polymers: non-oxide fibres via 68, 69‘pop-in’ stress 196pore diameter, effective 135–6

pore occlusion 290pore radius 134–5porosity 89, 90

and flexural strength 346–8fugitive additives and production of porous

ceramics 515–18silicon nitride and silicon carbide-based

ceramics 541thin films 232–5

porosity FGMs 577porous cell boundaries 15powder bed technique 170–3powder lamination 161–6powder mixing 286–7powder processing 155

particulate composites 100–3routes for FGMs 578–81

PRD–166 fibre 62–3preformed spinel-based refractories 365–6pressure-assisted sintering

WC-Co FGMs 589–90see also hot isostatic pressing (HIP); hot

pressing; spark plasma sintering(SPS)

pressureless sintering (PLS) 157, 341, 342, 343compared with HIP for carbon-ceramic

alloys 518–23FGMs 582–3

process zone 109processing

cermet nanocomposites 285–90damage due to 334FGMs 577–83fibre-reinforced glass/glass ceramic matrix

composites 72–80, 82–3fibrous monolithic ceramics 11–14infiltration processing of LGMs 137–8magnesia-spinel 362–3nanoceramics 255–6nanocomposite processing 514–15nanostructured silicon carbide coatings

273, 275–6, 277particulate composites 100–3sialon FGMs 161–74silicon nitride and boron carbide based

laminates 189–93whisker-reinforced silicon nitride ceramics

35–7processing conditions and R-curve

behaviour 47–8prompt gamma activation analysis (PGAA) 526pullout, whisker 41, 546–7, 554, 556–7pulsed electric current sintering (PECS) 273,

583Pyrex (borosilicate glass) 417

quasi-plasticity 143, 144, 146–7

Index 609

quasi-static model for erosion 552quench temperature 386–7

critical 372quench tests 407–8quenching temperature difference 407–8

critical 407–8, 411–12, 416–17, 425–7

R-curve behaviour 541, 555particulate composites 107, 109, 112, 117whisker-reinforced silicon nitride ceramics

46–8R parameter 373, 384–6, 410R´´´´ parameter 374, 384–6, 388–9, 390Rachinger grain boundary sliding 436Raj and Chyung model 442–3, 444reaction bonded silicon nitride (RBSN) 36–7reactive powder processing 583reduction 287–8refractive index 471, 473–4reinforced concretes 2relative strength 386–7residual strains 142–3residual stresses

alumina-lanthanum phosphate 341, 343–4FGMs

cracking 581–2thermal stress calculations 584–7

and fracture strength of cermetnanocomposites 299–300

hardening due to in thin films 230–2machinable nanocomposite ceramics 341,

352sialon composites 497–8

silicon carbide-reinforced 508–9and thermal shock of fibre-reinforced

CMCs 425–6residual stresses control 178–215

boron carbide-based laminates 201–9ballistic tests 203–4, 205, 208–9design and mechanical behaviour 201–3microstructures of three-layered boron

carbide/boron carbide-siliconcarbide laminates 206–9

undesirable influence of tensile residualstresses 204–6

future trends 210–11laminate design for enhanced fracture

toughness 179–89calculation of the residual stresses

179–81calculations of apparent fracture

toughness 185–7design principles and algorithm 181–5other toughening mechanisms 187–9

processing of silicon nitride-titaniumnitride and boron carbide-siliconcarbide laminates 189–93

silicon nitride-based laminates 193–201apparent fracture toughness 194–9fracture surfaces after fracture

toughness tests 199–201mechanical properties 193–4

retained strength after shock 388–9rice hulls 72rocksalt-structured compounds 235rolling 190, 191, 192rotary cement kilns 367–70, 391

problems in 368–70

saffil fibre 63, 65salt infiltration 289–90scanning electron microscopy (SEM) 92,

291–2non-oxide interfaces 466–7

scanning transmission electron microscopy(STEM) 463

Schottky barrier 326SCS fibres 68–9, 70seashells 1–2self-propagating high-temperature synthesis

(SHS) 304, 583semiconductors 221

CNT-ceramic composites 326–8, 330sequential slip casting 579shaping processes, for FGMs 578–81shear

along columnar grain boundaries 232–6shear-initiated failure 21–2, 23

shear bands 325shear modulus 222–6, 568shear stresses 219–22shear thickening creep 444–5sialon-based functionally graded materials

154–77functionally graded sialon ceramics 160–1production techniques 161–74

controlling the sintering conditions173–4

lamination technique 161–70powder bed technique 170–3

sialon composites 494–510a/b-sialon 155, 157, 160, 493, 494-8, 510carbon fibre-reinforced 500–5future trends 510particle-reinforced 498–9silicon carbide fibre-reinforced 505–9whisker-reinforced 499–500

sialons 155–60, 491–513a-b-sialon 155, 157, 160, 493, 494–8,

510a-sialon see a-sialonb-sialon see b-sialonchallenges in toughening and strengthening

493–4

Index610

dewetting behaviour 468–9microstructure 537–8, 556–7

eroded surfaces 544–5erosion performance 543–4and erosion mechanisms 553and mechanical properties 550–2

nitrogen and refractive index of 473–4O–sialon 155–6, 157, 159superplasticity 447

Sigma fibre (BP Sigma) 68, 70silica and alumina mix fibre 63, 64silicon carbide 462

carbon-ceramic alloys 519, 521, 522intergranular films 462, 466–7interphase boundaries

h-BN 471–2, 479, 480silicon nitride 475–9

material properties 336microstructure of silicon carbide-based

ceramics 536–59eroded surfaces 547–50erosion mechanisms 555–6erosion performance 543–4material characterisation 540–3mechanical properties 550–2material selection 539–40

monofilament-reinforced glass 82–3, 84, 85nanostructured coatings see nanostructured

coatingsNicalon/silicon carbide CMCs 421–3non-wetting behaviour 468oxynitride glass-silicon carbide composites

569–71particle-reinforced sialon composites 499processing silicon carbide fibres 68–71thermal shock behaviour 413thermal shock resistance parameters 411whisker-reinforced sialons 499–500whiskers 35

processing 72, 73silicon carbide/alumina nanocomposites

101–2, 103, 114–16, 118–19,122–4, 125, 244–5

silicon carbide/boron carbide laminates seeboron carbide-based laminates

silicon carbide fibre-reinforced sialons 505–9effects of heat treatment of fibres 506–7effects of sintering additives 507–8residual stresses 508–9

silicon carbide-reinforced silicon nitride 38–9,45–6, 49–52

silicon carbide-silicon nitride nanoceramics250–1

silicon carbide-titanium boride composites 117silicon carbide whisker-reinforced glass/glass

ceramic composites 87–92silicon carbonitride (HPZ) 71

silicon dioxide 323, 324silicon nitride 15, 323

calcium dopants in intergranular films474

carbon-containing see carbon-ceramicalloys

crystallographic modifications 155, 156fibres 71intergranular films 462, 466–7interphase boundaries

h-BN 479–82silicon carbide 475–9

material properties 336matrix 34–5microstructure of silicon nitride-based

ceramics 536–59eroded surfaces 546–7erosion mechanisms 553–5erosion performance 543–4material characterisation 540–3material selection 538–9mechanical properties 550–2

monolithic 16–17, 24–6need for sintering additives 156–7reinforcement of cell boundaries 15, 16silicon carbide-reinforced 38–9, 45–6,

49–52silicon nitride (w) reinforced silicon nitride

39–41wetting and non-wetting behaviour 468whisker-reinforced see whisker-reinforced

silicon nitride ceramicswhiskers 35

silicon nitride-based laminates 193–201apparent fracture toughness 194–9fracture surfaces after fracture toughness

tests 199–201, 202mechanical properties 193–4processing of silicon nitride-titanium

nitride laminates 189–93silicon nitride-boron nitride composites

fibrous monolithic ceramics 14mechanical properties 15–28

silicon nitride/h-BN machinablenanocomposites 344–51

silicon nitride-silicon carbide nanoceramics250–1

silicon/nitrogen mass ratio 527–8silicon oxide 366–7silicon oxynitride 564single-walled carbon nanotubes (SWCNTs)

251, 252–3, 310, 311, 312see also carbon nanotubes (CNTs)

sinter forging 449sintered alumina–iron CMC 414, 415sintered spinel clinker 363sintering

Index 611

controlling the sintering conditions forsialon FGMs 173–4

FGMs 582–3fibre-reinforced glass/glass ceramic

composites 73–4grain size reduction and superplasticity

445–6magnesia 363–4magnesia-alumina system 364magnesia-spinel system 365partial 515particulate composites 101–2, 103process and toughness of CNT

nanocomposites 252–3sialons 157–8see also under individual techniques

sintering additives 156–7, 515effect on silicon carbide fibre-reinforced

sialons 507–8, 509particulate composites 102, 103and superplasticity 447whisker-reinforced silicon nitride ceramics

34, 38–9, 40–1, 44, 50–1, 53–4sintering temperature

carbon fibre-reinforced sialons 500–1silicon-coated carbon nanotubes/WC-Co

composites 276sintering time

diffusion between layers of sialon FGMs164-5

and toughness of whisker-reinforcedsilicon nitride ceramics 49

slag resistance 371–2slurry infiltration process 74–5, 76soda lime glass 89–92soft or layered materials 335, 336soft nanoparticle dispersed nanocomposites

246–8sol-gel process

cermet nanocomposites 288–9fibre-reinforced glass/glass-ceramic matrix

composites 76–80, 95solid freeform fabrication (SFF) techniques

29solid state reactions 364solution-precipitation creep 442–4spalling 576spark plasma sintering (SPS) 304, 583

carbon nanotubes-ceramic composites 319,329–30

CNT-reinforced silicon nitride ceramics530–1, 532

nanoceramics 255, 256spatially resolved-valence electron energy loss

(SR-VEEL) spectroscopy 482spinels 359

formation by solid state reactions 364

magnesia-spinel ceramic composites seemagnesia-spinel ceramiccomposites

spontaneous microcracking 105, 106stable crack growth 197–9steady state erosion rate 543–4, 552step model 443–4stoichiometric magnesia-spinel 360

thermomechanical behaviour at hightemperatures 370–1

strain rate characterising superplasticity 437,439, 440, 441, 442, 444–5

straw and mud bricks 2strength

carbon fibre-reinforced sialon composites504, 505

cermet nanocomposites 299–300, 301challenges in strengthening sialons 493–4flexural see flexural strengthfracture strength see fracture strengthmagnesia-spinel ceramic composites

378–9hot strength 371, 376, 377and magnesia grain size 378–9relative strength 386–7retained strength after shock 388–9

particulate composites 110–20high-temperature 116–20room-temperature 110–16

retained after shock 388–9silicon nitride laminates 193–4silicon nitride and silicon carbide ceramics

536thermal shock in monolithic ceramics

412–13stress

distribution in a coating-substrate systemand a FGM 584

flow stress 218–19fracture stress test 46matrix cracking stress 426maximum tensile stress 404–6residual stresses see residual stresses;

residual stresses controlshear stresses 219–22and superplasticity 437–8thermal stresses see thermal stressesthreshold stress 197–9, 438

stress-induced microcracking 109stress intensity factor 182, 186–7, 188stress reduction factor 406stress-strain curves

alumina-matrix LGMs 143, 144YTZP multilayers 449, 450, 451, 452

sub-grain boundaries 245Sumitomo Chemical Company 63, 64superconductor ceramic matrices 328

Index612

superplastic ceramic composites 119–20,434–58

accommodation processes controllingsuperplasticity 439–45

GBS accommodated by diffusionalflow 439, 440

GBS accommodated by dislocationmovement 439, 440–2

shear thickening creep 444–5solution-precipitation creep 442–4

applications of superplasticity 448–52definition of superplasticity 434future trends 454macro- and microscopic superplastic

characteristics 435–9parameters improving superplasticity 445–8

improvement of processes controllingdiffusion 446–8

refinement of the microstructure 445–6properties to be controlled in

superplasticity 453superplasticity in nanoceramics 248–51

high temperature 250–1low temperature 248–50

surface activation 530surfactants 314

tailoring the architecture 15, 16tape lamination 166–70temperature

application of fibre-reinforced glass/glass-ceramic matrix composites 93–5

behaviour of stoichiometric spinel at hightemperatures 370–1

diffusion between layers of sialon FGMs164–5

flexural strength of whisker–reinforced siliconnitride ceramics and 43–4

high-temperature strength of particulatecomposites 116–20

oxidation of silicon carbide nanostructuredcoatings 270–1

properties of fibrous monolithic ceramics16–28

high temperature 21–8room temperature 16–20

sintering temperature 276, 500–1thermal expansion mismatch and Young’s

modulus at high temperatures formagnesia-spinel ceramiccomposites 370

tensile-initiated failure 21tensile modulus 85, 86tensile stresses

critical tensile stress 206maximum 404–6residual in boron carbide laminates 204–6

residual in silicon nitride laminates 194–9tensile tests 409tension: and compression in creep 437–9tetragonal zirconia polycrystalline materials

(TZP) 587–8tetrahedral sites 362Textron SCS-type silicon carbide fibres

68–9, 70thermal conductivity 43, 328–9thermal debond cracks 418–19, 420, 423–4thermal expansion mismatch 497–8

carbon fibre-reinforced sialon composites502–4

compensation in thermal mismatch504–5

FGMs 582magnesia-spinel composites 370sialon composites 497–8silicon carbide fibre-reinforced sialon

composites 508–9thermal shock of fibre-reinforced CMCs

425–6thermal microstresses 103–5

direct influence of 110thermal properties 43

advanced ceramics and layered or softmaterials 335, 336

CNT-ceramic composites 328–9, 330CNTs 312thermal shock damage and for fibre-

reinforced CMCs 417–25thermal shock 400–33

experimental methods 407–10assessment of thermal shock damage

409–10thermal shock simulation methods

407–8fibre-reinforced CMCs 416–27, 427–8

thermal shock damage and its effect onmechanical and thermal properties417–25

theoretical considerations 425–7induced stress field 401–7, 427magnesia-spinel ceramic composites

359–60, 372–4, 391–2mechanical properties and thermal

shock behaviour 375–89monolithic ceramics 400, 410–13particle- and whisker-reinforced CMCs

413–16resistance of fibrous monolithic ceramics

24–8resistance of whisker-reinforced silicon

nitride ceramics 52–3thermal shock damage resistance parameters

372, 374, 376, 377, 378, 384–6,388–9, 410–11

Index 613

thermal stress intensity factor 412thermal stress resistance parameters 372, 373,

384–6thermal stresses

calculation in FGMs 584–7carbon fibre–reinforced sialon composites

503–4induced stress field 401–7, 427rotary cement kilns 369

thicknessfilms at interfaces 463–7

see also equilibrium film thicknesslayer thickness and apparent fracture

toughness 182–3silicon carbide nanostructured coatings

263–4thickness ratio of layers 183–5thin films

deformation processes and microstructureof 232–6

residual stresses and hardening 230–2three-dimensional woven CMCs 421–43M 63three-phase microstructures 119threshold stress 197–9, 438titania 392titanium boride/aluminium multilayer 228, 229titanium boride-silicon carbide composites 117titanium carbide-alumina cutting tools 121, 124titanium-containing particles 122titanium nitride 216, 217, 222, 223, 224, 314–15

alumina-titanium nitride composite 108–9silicon nitride-titanium nitride laminates

189–93titanium nitride/aluminium nitride

multilayers 234–6titanium nitride/copper multilayer 228, 229titanium nitride/NbN multilayer 225–6titanium nitride/vanadium nitride

multilayer 219–20, 222titanium silicon carbide 335, 336, 337, 351titanium tetrachloride 136, 137Tonen SiNB 71top layers, residual stresses in 194–9tortuosity 133–4, 136toughening increment 110toughness 99–100

alumina-matrix LGMs 139–40, 149–50challenges in toughening sialons 493–4design of tough ceramic laminates by

residual stresses control see residualstresses control

fracture toughness see fracture toughnessand strength 110–13toughening mechanisms 10

microcrack toughening 104, 109, 112for nanoceramics 251–5

toughening particulate composites 105–10whisker-reinforced silicon nitride ceramics

41–2, 46–52transformation toughness 435transgranular cracks

induced by an intragranular particle 352,353

magnesia-spinel composites 382–3, 384sialons 545, 553

transition metal nitrides 216transition zone 162–5transmission electron microscopy (TEM) 92–3

accuracy for film thickness at interfaces463–7

Travitzky and Shlayen equation for infiltration133

tridymite (cordierite) 81triple junctions 290tungsten carbide see cemented carbideturbine blades 576two-dimensional woven CMCs 421–4two-pore-size model 135–6Tyranno 71

ultrasonic agitation 313unidirectional (UD) CMCs 417–19unstable crack growth 197–9

vacuum impregnation 77–8, 79vanadium nitride 216, 219–20, 222, 223, 226vapour-liquid-solid (VLS) process 72, 73Vickers indentation hardness test 47

alumina-matrix LGMs 143–4, 145, 146–7nanostructured coatings 274–5

viscosityglassy phase and superplasticity 441–2oxynitride glasses 570

volume diffusion 298volume fraction

carbon fibre-reinforced sialons 500–1elongated grains in a/b-sialon composites

496–7and thermal shock of fibre-reinforced

CMCs 425W-phase oxynitride glasses 567Wakai’s step model 443–4Washburn equation for infiltration 133WC/Co materials see cemented carbideweak interface phase 335weak nanoparticle dispersed nanocomposites

246–8wear

microstructure and 536–7, 543–50erosion performance 543–4erosion tests 541–3examination of eroded surfaces 544–50

particulate composites 120–4

Index614

resistance of whisker-reinforced siliconnitride ceramics 52–3

wear-resistant coating 151wear-resistant parts 94wet processing 578–81wet spraying 578, 579wetting behaviour 467–9whisker impingement 36–7whisker orientation 36, 48–9

and erosion performance 543, 544whisker pullout 41, 546–7, 554, 556–7whisker-reinforced composites 3–4

glass/glass-ceramic composites 87–92, 94sialons 499–500thermal shock 413–16


applications 54–5fabrication 35–7, 47–8materials 34–5

silicon nitride matrix 34–5whiskers 35

properties 37–54improved properties 41–54mechanical properties 43–54microstructure 38–41physical properties 42–3

whiskers, production of 71, 72, 73wollastonites 81, 564, 567wood 1work of fracture (WOF)

fibrous monolithic ceramics 16–17, 26, 27at high temperature 22–4, 25

magnesia-spinel ceramic composites 376,382–4, 389

woven CMCs 421–4wurtzite structure 235–6

X-phase sialon 156, 157X-ray diffraction (XRD) 93

Y-Mg-Si-Al-O-N glass-ceramic composites569–71

Y-PSZ 336Y-Si-Al-O-N glass-ceramic composites 561–2,

563–6Yajima process 68, 69Yokota equation for infiltration 133Young’s modulus 180–1, 183, 219, 542

alumina-lanthanum phosphate 339, 340boron carbide laminates 203carbon-ceramic alloys 528–9carbon fibre-reinforced sialon composites

505fibrous monolithic ceramics 18magnesia-spinel ceramic composites 376,

377, 379–80thermal expansion mismatch and 370

mismatch 222–6, 497–8oxynitride glasses 566, 568silicon nitride/h-BN machinable

nanocomposites 348silicon nitride laminates 193–4thermal shock and 423

fibre-reinforced CMCs 425whisker-reinforced silicon nitride ceramics 46

yttria 102, 103, 446sintering additive for whisker-reinforced

silicon nitride ceramics 34, 44,50–1, 53

yttria-alumina-zirconia (YAZ) sinteringadditive 507-8, 509

yttria-stabilised tetragonal zirconia polycrystalceramic (YTZP) 435, 437, 438,440, 441, 453

improvement of processes controllingdiffusion 446, 447–8

joining 449–52yttrium aluminium garnet (YAG) 118, 564yttrium disilicate 563–4, 569

zerodur 81zig-zag-type SWCNTs 310zirconia 80, 303-4, 392

alumina-zirconia/aluminium titanatesystems 140–7

alumina/zirconia FGM disks 580–1, 582,587, 588, 589

alumina/zirconia nanoceramics 253–5compensation for thermal mismatch in

carbon fibre-reinforced sialoncomposites 504–5

particles and sialon composites 499thermal shock resistance parameters 411transformation toughness 435

zirconia-based FGMs 587–8zirconia toughened alumina 254ZTA 138–40