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AerospaceMaterials

HandbookEdi ted by

Advances in Materials Science and Engineering

Series Editor

Sam Zhang

Aerospace Materials Handbook, edited by Sam Zhang and Dongliang Zhao

Biological and Biomedical Coatings Handbook: Applications, edited by Sam Zhang

Biological and Biomedical Coatings Handbook: Processing and Characterization, edited by Sam Zhang

CRC Press is an imprint of theTaylor & Francis Group, an informa business

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AerospaceMaterials

HandbookSAM ZHANG • DONGLIANG ZHAO

Edi ted by

Advances in Materials Science and Engineering

CRC PressTaylor & Francis Group6000 Broken Sound Parkway NW, Suite 300Boca Raton, FL 33487-2742

© 2013 by Taylor & Francis Group, LLCCRC Press is an imprint of Taylor & Francis Group, an Informa business

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vii© 2013 by Taylor & Francis Group, LLC

ContentsSeries Preface ....................................................................................................................................ixPreface...............................................................................................................................................xiEditors ............................................................................................................................................ xiiiContributors .....................................................................................................................................xv

Chapter 1 Superalloys for Super Jobs ...........................................................................................1

Seyid Fehmi Diltemiz and Sam Zhang

Chapter 2 Tool Condition Monitoring in Machining Superalloys ..............................................77

H. Z. Li and X. Q. Chen

Chapter 3 Laser Cladding and Alloying for Aerospace Applications ...................................... 109

S. K. Mishra

Chapter 4 High-Performance Wear/Corrosion-Resistant Superalloys ...................................... 151

Rong Liu and Matthew X. Yao

Chapter 5 High-Temperature Oxidation of Aerospace Materials ............................................. 237

Xiao Peng and Fuhui Wang

Chapter 6 Thermal Spray Coatings for Aeronautical and Aerospace Applications ................. 281

Chang-Jiu Li, Guan-Jun Yang, and Özge Altun

Chapter 7 Nanostructured Solid Lubricant Coatings for Aerospace Applications ................... 359

B. Deepthi and Harish C. Barshilia

Chapter 8 Processes and Characterizations of Metal Matrix Composites ............................... 415

Alokesh Pramanik and L. C. Zhang

Chapter 9 Processing Science for Polymeric Composites in Aerospace .................................. 461

Sunil C. Joshi

Chapter 10 Carbon Nanotube-Reinforced Polymer Composites for Aerospace Application ...... 493

Nanda Gopal Sahoo and Lin Li

For cataloging purposes only

viii Contents

© 2013 by Taylor & Francis Group, LLC

Chapter 11 Emerging Technology in Aerospace Engineering: Polymer-Based Self-Healing Materials ............................................................................................. 531

Jinglei Yang, He Zhang, and Mingxing Huang

Chapter 12 Preparation and Processing of Magnesium Alloys ..................................................607

Maoyin Wang, Yunchang Xin, and Qing Liu

Chapter 13 Fatigue of Magnesium Alloys .................................................................................. 647

F. A. Mirza and D. L. Chen

Chapter 14 Aerogel Materials for Aerospace ..............................................................................699

Yen-Lin Han


ix© 2013 by Taylor & Francis Group, LLC

Series Preface

ADVANCES IN MATERIALS SCIENCE AND ENGINEERING

SERIES STATEMENT

Materials form the foundation of technologies that govern our everyday life, from housing and house-hold appliances to handheld phones, drug delivery systems, airplanes, and satellites. Development of new and increasingly tailored materials is key to further advancing important applications with the potential to dramatically enhance and enrich our experiences.

The Advances in Materials Science and Engineering series by CRC Press/Taylor & Francis Group is designed to help meet new and exciting challenges in materials science and engineering disci-plines. The books and monographs in the series are based on cuttingedge research and development, and thus are up-to-date with new discoveries, new understanding, and new insights in all aspects of materials development, including processing and characterization and applications in metallurgy, bulk or surface engineering, interfaces, thin �lms, coatings, and composites, just to name a few.

The series aims at delivering an authoritative information source to readers in academia, research institutes, and industry. The publisher and its series editor are fully aware of the importance of materials science and engineering as the foundation for many other disciplines of knowledge. As such, the team is committed to making this series the most comprehensive and accurate literary source to serve the whole materials world and the associated �elds.

As series editor, I would like to thank all authors and editors of the books in this series for their noble contributions to the advancement of materials science and engineering and to the advance-ment of humankind.

Sam Zhang


xi© 2013 by Taylor & Francis Group, LLC

Preface

Flying a machine into the sky requires tremendous thrust and power to launch and sustain the �ight, be it an airplane or a space shuttle. The engine thus needs “superalloys” to function properly at extreme temperatures as high as over one thousand degrees Celsius; and the body of the �y-ing machine needs lightweight yet strong materials thus the need for new materials arises such as polymer composites, magnesium alloys, and so on. Materials used for aeronautical and aerospace applications are perhaps one of most advanced materials of any time, owing to its harsh applica-tion environment and the most stringent safety requirements. Since the Wright Brothers over a 100 years ago, the aviation landscape has changed tremendously, and so have the materials used to make that happen. With the development of superalloys that work at even higher temperatures with even lighter weight, we can expect faster and more powerful airplanes for passenger, cargo, and all air or space applications.

This book covers traditional superalloys and recent development and applications of light alloys such as magnesium alloys and aerogel materials. It provides a timely handbook for seasoned researchers as well as newcomers in the aerospace materials �eld.

This book comprises chapters dealing with bulk materials, coatings, and traditional and new materials for aerospace applications, from preparation/processing to machining, and from oxidation/corrosion to coating protection. There are 14 chapters in the book. Chapter 1, Superalloys for Super Jobs, overviews the �eld of superalloys concentrating on traditional nickel−iron-based superalloys, nickel-based superalloys, and cobalt-based superalloys. Chapter 2 focuses on Machining and Chapter 3 on Laser Cladding and Alloying. The Wear/Corrosion Performance of superalloys is summarized in Chapter 4, while Chapter 5 pays special attention to High Temperature Oxidation. Chapter 6 elaborates on coating protection through Thermal Spraying and Chapter 7 introduces Nanostructured Solid Lubricant coatings. Chapters 8 through 11 cover four categories of com-posites used in aerospace: Metal Matrix Composites, Polymer Composites, Carbon Nanotube Reinforced Polymer Composites, and Self-Healing Composites. Chapters 12 and 13 deal with the lightweight magnesium alloys: Preparation and Processing, and Fatigue. Aerogel is a synthetic porous material derived from a gel, in which the liquid component of the gel has been replaced with a gas. The result is a solid with extremely low density and thermal conductivity. Aerogel is a new class of materials that increasingly �nds exciting applications in aerospace and aeronautical applica-tions. Aerogel Materials is introduced at the end of book in Chapter 14.

This handbook is different from traditional the “handbook” because it is compiled not just with data (tables and �gures). It is written in such a way that both novices and veterans will �nd it useful. This book aims to bring up the reader’s knowledge horizon to the forefront of the research in materi-als for aerospace and aeronautical applications. For newcomers, it will serve as a stepping stone into this �eld of materials; for veterans, the tables, �gures, and research highlights will be valuable to assist or stimulate their state-of-the-art research. Researchers in the materials �elds, materials scien-tists, engineers, and postgraduate students, especially those dealing with aerospace or aeronautical studies, should �nd the book useful.

The idea of the book was conceived and developed during Sam’s research attachment in the Central Iron and Steel Research Institute (CISRI, where Dongliang works), Beijing, China. CISRI plays a leading role in China’s R&D in superalloys. The editors would like to thank all the CISRI support and thank all the contributing authors for their painstaking work that �nally resulted in this nice book. Special thanks go to the reviewers of the chapters who patiently read through about 100 manuscript pages for each chapter (1 or 2 reviewers for each chapter) to provide their professional


xii Preface


critique that enabled the chapters’ improvement to their present form. Their professionalism and the contributions were the guarantee of the quality of the book. Last, but not least, the editors would also like to thank CRC Press staff, especially project managers Allison Shatkin and Jennifer Ahringer at Taylor & Francis Group, for their invaluable assistance rendered throughout the entire endeavor that made the smooth publication of the book a reality.

Sam ZhangProfessor, Nanyang Technological University

Singapore

Dongliang ZhaoProfessor, Central Iron & Steel Research Institute

People’s Republic of China


xiii© 2013 by Taylor & Francis Group, LLC

Editors

Professor Sam Zhang Shanyong, better known as Sam Zhang, received his PhD in ceramics in 1991 from The University of Wisconsin-Madison, USA and is a tenured full profes-sor at the School of Mechanical and Aerospace Engineering, Nanyang Technological University, Nanyang, China. Professor Zhang serves as the editor-in-chief for Nanoscience and Nanotechnology Letters and the principal editor for Journal of Materials Research (USA).

Professor Zhang’s association with superalloys dates back to his college years. In 1977, when China ended the infamous 10 years internal turmoil and resumed the normalities of the universities, Professor Zhang pass the nationwide university entrance examination and entered the “Class of 77 Superalloys”

of the Northeastern University, Shenyang, China. His bachelor’s thesis was on recrystallization of the GH28 superalloy, and his master’s thesis was on the hot-isostatic pressing densi�cation mechanisms of the K17 cast superalloy, obtained at the Central Iron and Steel Research Institute, Beijing, China.

After his doctoral studies on ceramics at UW-Madison, Professor Zhang entered the �eld of cermets and engineering coatings/thin �lms. He has been involved in the processing and charac-terization of nanocomposite thin �lms and coatings for 20 years, has authored and co-authored more than 260 peer-reviewed international journal papers with an average of more than 12 cita-tions per paper, 7 books, 20 book chapters, and has guest-edited more than 20 journal volumes. His book Materials Characterization Techniques (by Sam Zhang, Lin Li, and Ashok Kumar, published by CRC Press, USA, 344pp., publ. date: December 22, 2008, ISBN: 9781420042948; 1420042947) has been adopted as a textbook by eight American universities: Purdue University, New York University, Rutgers University, Johns Hopkins University, North Seattle Community College, Louisiana State University, California Polytechnic State University, and University of Missouri, and one European university: University of Southern Denmark. This book was also trans-lated into Chinese and published by China Science Press in October 2010. His other books are: Nanocomposite Thin Films and Coatings: Processing, Properties and Performance, edited by Sam Zhang and Nasar Ali, published by Imperial College Press, 628pp., publ. date: October 2007, ISBN: 978-1-86094-784-1; 1-86094-784-0; a three-volume set edited by Sam Zhang, published by CRC Press Taylor & Francis Group on June 18, 2010: Vol. 1, Nanostructured Thin Films and Coatings: Mechanical Properties, ISBN: 9781420094022; ISBN 10: 1420094025; hardback, 550pp.; Vol. 2, Nanostructured Thin Films and Coatings: Functional Properties, ISBN: 9781420093957; hard-back, 422pp.; Vol. 3, Organic Nanostructured Thin Film Devices and Coatings for Clean Energy, ISBN: 9781420093933; hardback, 254pp.; and a two-volume set edited by Sam Zhang, published by CRC Press Taylor & Francis Group in May 2011: Vol. 1, Biological and Biomedical Coatings Handbook: Processing and Characterization, ISBN: 978-1-4398-4995-8; Vol. 2, Biological and Biomedical Coatings Handbook: Applications, edited by Sam Zhang, ISBN: 978-1-4398-4996-5. In addition, two more books are in in press: Hydroxyapatite Coatings (ed. Sam Zhang, CRC Press, 2013) and Surface Technologies (ed. Guojun Qi and Sam Zhang, Springer, 2014).

Professor Zhang was conferred the title of honorary professor of the Institute of Solid State Physics, Chinese Academy of Sciences. He also holds guest professorship at Shenyang University, Central Iron and Steel Research Institute, Zhejiang University, and Harbin Institute of Technology.


xiv Editors


Professor Zhang was featured in the �rst ever Who’s Who in Engineering Singapore (2007) and was also featured in the 26th and 27th editions of Who’s Who in the World. He became a Fellow of the Institute of Materials, Minerals and Mining, UK, in October 2007. He has been called to pres-ent over 60 invited or plenary keynote lectures at international conferences in Taiwan, Japan, the United States, France, Spain, Mainland China, Portugal, New Zealand, and Germany. He has also been invited over 20 times by universities and industries to conduct short courses and workshops. He founded the Thin Films conference series in 2002 and has been the chair of this very successful conference series ever since. Professor Zhang is also the president of the Thin Films Society estab-lished in Singapore in 2009.

Professor Zhang’s current research centers on the following four aspects: 1. Hard yet tough nanocomposite coatings for tribological applications by physical vapor deposition; measurement of fracture toughness of ceramic �lms and coatings. 2. Biological coatings and drug delivery applica-tions. 3. Electronic thin �lms in optoelectronic applications, and data storage. 4. Energy �lms and coatings for dye-sensitized solar cells. Details of Professor Zhang’s research are accessible at http://www.ntu.edu.sg/home/msyzhang

Professor Dongliang Zhao received his PhD in 1993 from the Central Iron and Steel Research Institute (CISRI), Beijing, China, on microstructure studies of superalloys. He has been the chief engineer at CISRI since 2009 and the director of CISRI’s Institute of Functional Materials. Professor Zhao’s research interests include computational material science, magnetic materials, and energy materials as well as superal-loys. Professor Zhao has been the leading principal investigator of, or participated in, more than 20 Chinese national research projects funded by the National Natural Science Foundation of China (NSFC) program, the National Key Basic Research program, the National High Tech Research and Development Program of China of the General Armament Department, and so on. These projects include “New-generation single crys-

tal superalloy interface structure and solid solution effects” from NSFC, “Aero-engine superal-loy microstructure and macroscopic properties” from the General Armament Department, among others, and participated in superalloy design projects funded by the National Key Basic Research program. Apart from classi�ed research, Professor Zhao has published some 40 journal papers and has been granted six patents.

In 2003, Professor Zhao was conferred the title of “Beijing Outstanding Young Engineer” by the Beijing City Government. In 2006, he was recognized by the State Department as one of the National Star Researchers, and in 2008, he was conferred the title of “National Defense Science and Technology Innovation Leader.”


xv© 2013 by Taylor & Francis Group, LLC

Contributors

Özge AltunMechanical Engineering DepartmentEskisehir Osmangazi UniversityEskisehir, Turkey

Harish C. BarshiliaSurface Engineering DivisionCSIR National Aerospace LaboratoriesBangalore, India

D. L. ChenDepartment of Mechanical and Industrial

EngineeringRyerson UniversityToronto, Ontario, Canada

X. Q. ChenDepartment of Mechanical EngineeringThe University of CanterburyChristchurch, New Zealand

B. DeepthiSurface Engineering DivisionCSIR National Aerospace LaboratoriesBangalore, India

Seyid Fehmi DiltemizAerospace Material LaboratoryAircraft Maintenance CenterEskisehir, Turkey

Yen-Lin HanDepartment of Mechanical EngineeringUniversity of ConnecticutStorrs, Connecticut

Mingxing HuangSchool of Mechanical and Aerospace

EngineeringNanyang Technological UniversitySingapore

Sunil C. JoshiSchool of Mechanical and Aerospace


Chang-Jiu LiSchool of Materials Science and

EngineeringXi’an Jiaotong UniversityXi’an, People’s Republic of China

H. Z. LiSchool of Mechanical and Manufacturing

EngineeringThe University of New South WalesSydney, Australia

Lin LiSchool of Mechanical and Aerospace


Qing LiuCollege of Materials Science and

EngineeringChongqing UniversityChongqing, People’s Republic of China

Rong LiuDepartment of Mechanical and Aerospace

EngineeringCarleton UniversityOttawa, Ontario, Canada

F. A. MirzaDepartment of Mechanical and Industrial

EngineeringRyerson UniversityToronto, Ontario, Canada


xvi Contributors


S. K. MishraCSIR-National Metallurgical LaboratoryCouncil of Scienti�c & Industrial ResearchJamshedpur, India

Alokesh PramanikDepartment of Mechanical EngineeringCurtin UniversityBentley, Australia

Xiao PengInstitute of Metal ResearchShenyang, People’s Republic of China

Nanda Gopal SahooSchool of Mechanical and Aerospace


Fuhui WangInstitute of Metal ResearchShenyang, People’s Republic of China

Maoyin WangCollege of Materials Science and EngineeringChongqing UniversityChongqing, People’s Republic of China

Yunchang XinCollege of Materials Science and EngineeringChongqing UniversityChongqing, People’s Republic of China

Guan-Jun YangSchool of Materials Science and

EngineeringXi’an Jiaotong UniversityXi’an, People’s Republic of China

Jinglei YangSchool of Mechanical and Aerospace


Matthew X. YaoKennametal StelliteBelleville, Ontario, Canada

He ZhangSchool of Mechanical and Aerospace


L. C. ZhangSchool of Mechanical and Manufacturing

EngineeringThe University of New South WalesSydney, Australia

Sam ZhangSchool of Mechanical and Aerospace



1© 2013 by Taylor & Francis Group, LLC

Superalloys for Super Jobs

Seyid Fehmi Diltemiz and Sam Zhang

1CONTENTS

1.1 Introduction ..............................................................................................................................21.2 Historical Background ..............................................................................................................21.3 Types of Superalloys .................................................................................................................5

1.3.1 Strength and Strengthening Mechanisms of Superalloys .............................................51.3.1.1 Solid Solution Strengthening .........................................................................61.3.1.2 Precipitation Strengthening ...........................................................................81.3.1.3 Effect of Grain-Related Properties on Strengthening.................................. 151.3.1.4 Effect of Alloy Chemistry ........................................................................... 16

1.3.2 Iron–Nickel-Based Superalloys .................................................................................. 161.3.3 Nickel-Based Superalloys ........................................................................................... 181.3.4 Cobalt-Based Alloys ...................................................................................................22

1.4 Production and Processing Techniques for Superalloys .........................................................241.4.1 Casting–Melting Practice ...........................................................................................36

1.4.1.1 Melt Processing ............................................................................................361.4.1.2 Investment Casting .......................................................................................40

1.4.2 Wrought Alloys ........................................................................................................... 451.4.2.1 Mill Products: Primary Hot Working ..........................................................461.4.2.2 Forging .........................................................................................................481.4.2.3 Forming ........................................................................................................ 49

1.4.3 Powder Metallurgy ..................................................................................................... 511.4.3.1 Powder Consolidation Techniques ...............................................................54

1.4.4 Welding ....................................................................................................................... 551.4.4.1 Gas Tungsten Arc Welding and Gas Metal Arc Welding ............................ 571.4.4.2 Resistance Welding ...................................................................................... 581.4.4.3 Electron Beam Welding ............................................................................... 581.4.4.4 Laser Welding .............................................................................................. 591.4.4.5 Friction Welding .......................................................................................... 591.4.4.6 Conclusion ....................................................................................................60

1.4.5 Brazing ........................................................................................................................601.4.6 Machining ................................................................................................................... 62

1.4.6.1 Conventional Machining .............................................................................. 621.4.6.2 Water Jet ....................................................................................................... 631.4.6.3 Electrodischarge Machining ........................................................................ 631.4.6.4 Electrochemical Machining .........................................................................641.4.6.5 Laser Machining ..........................................................................................65

1.4.7 Heat Treatment of Superalloys ...................................................................................651.5 Main Failure Mechanisms of Superalloys ..............................................................................68

1.5.1 High-Temperature Oxidation ......................................................................................681.5.2 Hot Corrosion .............................................................................................................69


2 Aerospace Materials Handbook


1.1 INTRODUCTION

The term “superalloys” means alloys that are superior in heat and corrosion resistance, and main-tain superior properties even at elevated temperatures. Thus, superalloys are synonymous with “high-temperature alloys.” Traditionally, superalloys are basically classi�ed into three types according to their base element: iron-based superalloys, nickel-based superalloys, and cobalt-based superalloys. Presently, however, in a more broad sense, alloys that have the capability of preserving their mechanical, physical, and chemical stabilities at high temperatures and in severe corrosive environ-ments are all called superalloys.

At room temperature, however, the mechanical properties of superalloys are not much different from that of steel, which are much cheaper and easily produced. However, superalloys stand out from other metallic alloys with their high corrosion resistance even in this temperature range, making them ideal candidates for use in severe corrosive environments such as offshore petroleum platforms.

The production of superalloys is expensive. Not surprisingly, superalloys main application area is at elevated temperatures despite impressive room temperature and cryogenic properties. At high temperatures, superalloys are unique with their mechanical and chemical properties. The big leap forward for superalloys was a result of intense effort in developing gas turbine hot section parts to meet the needs of military aircraft. Today, both military and civil aircraft gas turbines still play a dominant role in the development of superalloys. Aviation gas turbines take up to 3/4 of all applica-tions of superalloys, the other 1/4 goes to power-generation gas turbines, chemical industry, medical industry, petrochemical equipment, space ships, rocket engines, nuclear reactors, submarines, high-temperature industrial furnaces, and various other applications that need high temperature and/or chemical resistance (Stoloff 1990).

Today, the energy sector strategically plays an even more critical role than ever before because of escalating fuel price. Global warming and increasing environmental pollution further necessitate more ef�cient use of petroleum products. The mandate of decreasing fuel consumption and exhaust emissions for aviation and industrial gas turbines is clearly paramount. This dictates the key posi-tioning and new development direction of all superalloys.

1.2 HISTORICAL BACKGROUND

Although superalloys owe their developments to aircraft gas turbines, their existence is older than aircraft gas turbines. Steam turbines began their journey in the 1800s and gas turbines were used in the 1900s for power generation. In 1917, the Imphy company in France took a patent for an alloy ATG (alliage pour turbines à gaz), which was developed for land-based gas turbines. Beside Fe, this alloy contained 60% Ni, 11% Cr, and was hardened by W and C. Chevenard at Imphy took a second patent in the same year for an alloy ATV (alliage pour tur-bines à vapeur), which was developed for the solution of intergranular corrosion problems in steam turbine blades. ATV was an austenitic iron-based alloy, which contained 34% Ni, 11% Cr, and was hardened by 0.3% C. After realizing that chromium carbide itself can cause corrosion problems, Chevenard thought of utilizing another strengthening mechanism known as precipita-tion hardening. He discovered the hardening effect of aluminum and titanium on these alloys, then patented it in 1929 (Durand-Charre 1997). Six months earlier (but published 6 years later), USA Inco also took a very similar patent for the effect of titanium and aluminum on well-known 80-20 Ni–Cr electrical heaters. Wilhelm Rohn produced Ni–Fe–Cr heat and corrosion-resistant alloys in Germany using vacuum induction melting (VIM) technology in the 1920s and the early

1.5.3 Fatigue ........................................................................................................................ 701.5.4 Creep ........................................................................................................................... 72

1.6 Recent Advances in Superalloys ............................................................................................. 74References ........................................................................................................................................ 74


3Superalloys for Super Jobs


1930s. The invention of aluminum and titanium precipitation hardening effect prepared suitable conditions for the development of well-known Nimonic 80 alloys by Pfeil et al. in England in 1941. Carbide-strengthened cobalt-based alloys were also seen in this era with the advantage of relatively easier castability. Only several years later, Batteridge and Franklin used x-ray dif-fraction techniques and detected the γ ′ Ni3(Al,Ti) phase in 1957 (Durand-Charre 1997) and thus solved the mystery of the aluminum and titanium precipitation hardening effect.

Superalloys were developed in parallel to the aircraft gas turbines. The �rst practical aircraft gas turbines were used in the 1937 �ight of Hans von Ohain’s Heinkel in Germany and in the 1939 �ight of Whittles’ engine in England. Many scienti�c experiments were performed to develop high-temperature alloys during World War II and these efforts built the basis of new alloys.

The 1950s and the 1960s saw the development of many new alloys by virtue of the effort dur-ing World War II. These alloys were usually used as-cast due to ease in production. However, the lack of knowledge in heat treatment and impurities caused large differences in properties among production lots. The development of superalloys had been closely related to inventions in other areas; for instance, the dramatic effect of boron in superalloys was only understood after developments in analytical chemistry, which allowed the determination of the amount of boron with higher sensitivity. Similarly, the development of the electron microscope made observation of the γ ′ phase possible. The usage of vacuum metallurgy on an industrial scale was also possible at the beginning of the 1950s. The value of vacuum melting was price-less; �rst, the easily oxidized reactive elements such as Al and Ti could be used at higher amount; also, melting time could be longer, which allowed the adjusting of alloy composition with higher accuracy; and alloy cleanliness was greatly improved by longer melting time and evaporation of volatile impurities, such as Bi, under vacuum. Nevertheless, the 1950s and the 1960s were the years of new alloys and the 1970s and the 1980s were the years of process devel-opments (Tien and Caul�eld 1989).

It will never be a complete history of superalloys without mentioning the coatings. The �rst-generation coatings were developed in the 1950s and the 1960s aiming at the protection of alloys from oxidation and hot corrosion damages. These coatings were produced by diffusion of alumi-num using pack cementation technique. By virtue of these coatings, the alloy designers felt free in designing new alloys without corrosion damage. Successively, in the following decades, coatings were developed not only against corrosive environment but also against excessive heat. Thermal barrier coatings (TBCs) were �rst produced by plasma spray, and then by electron beam physical vapor deposition (EB-PVDs) (Diltemiz 2010). With TBCs, the surface temperature of the alloy decreased up to 200°C under excessive heat transfer conditions; thus, more ef�cient gas turbines with higher turbine inlet temperature (TIT) became possible. Today, the widely used TBC struc-ture consists of two layers: MCrAlY bond coat and 7% Y2O3-stabilized zirconia top coat (Portinha et al. 2005.)

The inventions in powder metallurgy and mechanical alloying made the production of oxide dispersion strengthening (ODS) alloys possible. New techniques such as rapid solidi�cation, super-plastic forming, and near net shape process signi�cantly increased the process ef�ciency of superal-loys. Directional solidi�cation (DS) and its logical extension single crystal (SX) technology induced the birth of a series of new superalloys since the 1980s. Thanks to the lack of grain boundaries, SX alloys are free from grain boundary-related weaknesses. Grain boundary strengthener elements such as C cause decreasing incipient melting point of the alloys. SX allows the arrangement of new alloy compositions without grain boundary elements. Newer series of superalloys started using rhenium and ruthenium elements. Some of the major developments in superalloys are given chrono-logically in Figure 1.1.

Today, increasing reliability of coatings, production techniques, and new material development efforts still continue. In addition, intense research and development work is ongoing in the so-called fourth generation of SX superalloys with less expensive rhenium and ruthenium elements (Reed 2006).



2010

2000

1990

1980

1970

1960

1950

1930

1920

1910

1900Invention of vacuum

induction melting (VIM) byColby, the USA

Imphy’s ATG and ATV iron-basedheat and corrosion resistant alloys.

France

Wilhelm Rohn, vacuuminduction melted nickel–iron–

chromium alloys, Germany

Realizing of A1 and Tielements’ hardening

effect. France and the USA

Frank Whittle’sengine

England

Vacuum induction melting(VIM) application insuperalloys, waspaloy

1952, the USA

Development of selfsupporting shells ininvestment casting

Development ofdirectionally solidified

(DS) alloys

Adoption in HIPtechnology to thecast superalloys toeliminate porosity

Second generationSX superalloys with

using of rhenium

Process related

Alloy related

Otherdevelopments

1940

Fourth generation SXsuperalloys with

addition of ruthenium

Third generation SXsuperalloys with usingof higher amount of

rheniumEB-PVD thermal

barrier coatings (TBCs)

Plasma sprayedthermal barriercoatings (TBCs)

Invention ofmechanical

alloying (MA)

Pack aluminidecoatings application

of superalloys

First γ' phasedetermination by usingx-ray diffraction, 1957

Invention ofimportance of boron

PFEİL et al. firstprecipitation harnedednickel-based superalloynimonic 80, England

Heinkel's iron-based tiniduralloys, Germany

Iron–nickel-basedA-286

Co-based vitallium andstellite alloys, the USA

First γ ' phaseobservation byusing scanning

electronmicroscope (SEM)

Application ofpowder metallurgyand MA techniquefor production ofODS superalloyssuch as MA 754,

MA 758 etc.

Development ofsingle crystal (SX)

superalloys

CMSX-4

CMSX-10TMS-75

TMS 162

Rene N5

Rene 95

Udimet 720

SRR99MA 6000

MAR-M 509MAR-M 200

X-40

Hastelloy B

Udimet 700

Inconel 100

Rene N6

CMSX-2

Eiselstein's iron–nickel-based alloy

inconel 718

FIGURE 1.1 Some of the major developments in superalloys.




1.3 TYPES OF SUPERALLOYS

As seen in the historical evolution, mainly three types of superalloys exist, the �rst being iron–nickel-based alloy. Like all other engineering materials, the selection of superalloys pulls together many factors such as application conditions, production feasibility, price, and so on. The vast range of variations in service conditions caused the development of several alloys with very different properties.

Gas turbine ef�ciency depends on TIT. Therefore, higher TIT demands improvement in many critical design parameters such as reduced fuel consumption, increased range, reduced exhaust emission, increased load-carrying capacity, and so on. Thus, gas turbine designs have successively aimed at higher TIT; and alloy designers have developed new alloys of higher temperature resis-tance. The typical relationship between TIT and fuel consumption is shown in Figure 1.2.

1.3.1 STRENGTH AND STRENGTHENING MECHANISMS OF SUPERALLOYS

The main strength requirement of superalloys is considerably different from conventional engi-neering alloys. The strength of conventional metallic alloys usually implies tensile strength, yield strength, hardness, and ductility. The main strength requirements of superalloys are de�ned below:

Tensile properties: For superalloys, the tensile properties are always spoken of with respect to temperature. Related properties include yield and tensile strength, ductility, and elastic-ity (stiffness). One of the special forms of tensile property important for superalloys is the stress-rupture strength (creep-rupture strength), or the failure of component under static load at elevated temperatures and speci�ed environment.

Creep resistance: Creep is a permanent plastic deformation with time under high tempera-ture and loads normally well below the yield point. Creep damage can go up to complete separation and their detrimental effect is irreversible. Under stress and high temperature, gas turbine parts such as disks, bolts, and blades are subjected to creep damage; therefore, creep resistance is a critical parameter for turbine designers. Three parameters are needed for creep damage to occur: time, temperature, and stress. Slow diffusing elements improve creep resistance of superalloys. Creep damage usually occurs in grain boundaries espe-cially in triple points (junction point of the three grains); thus, grain boundary properties are also important.

Fatigue resistance: Fatigue is another important failure type causing crack propagation or complete rupture of components under cyclic stress below ultimate strength. The cyclic

0500 700 900 1100 1300 1500 1700 1900

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10

Spec

ific

fuel

con

sum

ptio

n(g

r/kw

h)

Turbine inlet temperature (TIT) (°C)

FIGURE 1.2 Typical relationship between TIT and fuel consumption. (Adapted from Diltemiz S.F. 2010. Thermal and mechanical properties optimization of thermal barrier coatings. PhD dissertation, ESOGU University: Turkey.)



stresses may not only be mechanical but also thermal or mechanical and thermal at the same time usually for aerospace component, resulting in thermomechanical fatigue instead of simple thermal fatigue.

Corrosion resistance: Corrosion resistance re�ects a material’s ability to resist chemical deg-radation in a speci�c corrosive environment. Corrosion becomes a complex phenomenon in superalloys. Many corrosion-related failure mechanisms exist in superalloys, including stress corrosion cracking (SCC), intergranular oxidation (IGO), intergranular attack (IGA), high-temperature oxidation, stress-assisted grain boundary oxidation (SAGBO), hot cor-rosion, and so on.

Resistance against these damage mechanisms depends on alloying chemistry, microstructural design, and heat treatment history, and production technique. All of these affect the alloy’s strength (sometimes negatively) and require careful balance between properties and application conditions for a particular alloy. Above all, the main strength-limiting factors of a superalloy are incipient melting and dissolution of the useful phases under service temperature.

There are three principal strengthening mechanisms for superalloys: solid solution strength-ening, precipitation hardening of intermetallics, and carbide precipitation (Askeland 1985). Besides these mechanisms, controlling of crystal orientation, grain boundary size, grain aspect ratio (GAR), and orientation, improved purity by reducing the amount of trace elements, con-trolling microstructure by heat treatment, and careful application of the production process add further strength (Decker 1979). The principal strengthening mechanisms of superalloys are given in Table 1.1.

Superalloys have austenitic face-centered cubic (FCC) matrix called gamma (γ) phase. This phase is very ductile and suitable for solid solution and dispersion strengthening. Iron has body- centered cubic (BCC) and cobalt has hexagonal close-packed (HCP) structure at room temperatures. Both elements undergo allotropic phase transformations to FCC at high tempera-tures. Iron- and cobalt-based superalloys stabilize their FCC structures by means of austen-ite stabilizer alloying elements. Because of higher atomic mobility, FCC lattice structure has better creep resistance than BCC structure. In addition, FCC structure has higher high-tem-perature strength than BCC because atoms are more densely packed in the FCC structure. To obtain creep and oxidation resistance, it is necessary to add solution hardening elements into the matrix to form a solid solution. FCC matrix has much more solubility of hardening elements than BCC to atoms like Mo, Co, and W. Figure 1.3 illustrates rupture strength under constant temperature for 100 h.

1.3.1.1 Solid Solution StrengtheningSuperalloy matrices all have some solubility to alloying elements. These alloying elements are introduced as substitution atoms to the original lattice and distort it due to atomic radii difference. This situation is illustrated in Figure 1.4. Therefore, solid solutioning strength is achieved by lattice distortion that produces strain area serving as obstacles to the dislocation movement.

Solid solution strengthening decreases the stacking fault energy of the FCC lattice. This energy change is particularly important for inhibiting dislocation cross slip (which is the main deformation mechanism at high temperature). Solid solutioning increases room temperature yield and tensile strength of superalloys to some degree but critical improvement is achieved at high temperatures. Unlike dispersion hardening, solid solution matrix does not undergo cata-strophic change at elevated temperatures; thus, creep resistance is improved and loss of strength is hindered. Although solid solution strengthening increases the yield strength, tensile strength, and hardness, ductility, which is a desired critical property of superalloys, is reduced. The effec-tiveness of solid solution strengthening in superalloys is directly proportional to the atomic diameter difference, which controls the lattice distortion and the amount of alloying elements (Brooks 1982). As atomic diameter difference increases, more strength is obtained. Atomic




TABLE 1.1Principal Strengthening Mechanisms of Superalloys

Properties Comments

Solid solution strengthening

Its effectiveness depends on the atomic size difference between solute and solvent and will be able to stabilize to elevated temperatures. Therefore, although its ef�ciency on mechanical strength is limited, very useful to improve creep and stress rupture resistance. Mo, Co, and W are important solid solution hardener elements. Solubility in face-centered cubic (FCC) matrix is higher than that in body-centered cubic (BCC). Fe- and Co-based alloys are alloyed with Ni solid solution to gain FCC matrix structure.

Precipitation hardening of intermetallics

The main strengthening mechanism of nickel- and iron-based superalloys. The most popular and �rst discovered is γ ′ (Ni3Al, Ti, Ta, Re). While γ ′ can be used in both Ni- and Fe-based alloys, γ ′′ (Ni3Nb) is unique to iron-based alloys. γ ′ can retain their stability up about to 1150°C, which is well above the γ ′′ (only about 650°C).

Precipitation hardening of carbides

The main strengthening mechanism of cobalt-based superalloys. Therefore, the C content of cobalt-based superalloys is fairly higher than other groups. The brittleness and reduced feasibility of production are the main disadvantages of carbide strengthening.

Grain size Important and usually adjustable parameter. Its behavior depends on the temperature and component thickness. Large grains mean less grain boundary in the alloys. Usually large grains in gas turbine parts are desirable.

Grain aspect ratio (GAR)

The ratio of grains’ long axis to short axis, especially important in MA (mechanically alloyed) ODS (oxide dispersion strengthening) series superalloys.

Grain orientation Very useful strengthening mechanism when applied loads are anisotropic such as turbine blades. The desired orientation can be obtained by directional solidi�cation and rolling technique for cast and wrought alloys, respectively.

Crystallographic orientation

Some orientations are preferable due to higher strength in FCC matrix, such as <100>. It can be obtained by SX (single crystal) casting technique.

Alloy cleanliness Cleanliness can dramatically alter the properties of the alloys. It usually causes drop in incipient melting temperature and unwanted phases formation, which causes brittleness.

800

Carbide-hardened cobalt-based superalloys

Solid solutionsuperalloys

Precipitation-hardened nickeland iron-based superalloys

700

600

500

Stre

ss (M

Pa)

400

300

100

0700 800 900 1000

Temperature (°C)1100 1200 1300600

200

FIGURE 1.3 Stress rupture curves of superalloys. (Reprinted with permission of Davis, F.R. 1997. ASM Specialty Handbook Heat Resistant Materials, p. 221, ASM International: USA.)



diameter differences of some solid solution hardener in nickel-based superalloys are given in Table 1.2.

Aluminum is known as a perfect precipitation hardener but also as a potent solid solution hardener. Slow diffusing elements such as molybdenum and tungsten are very effective solid solution hardeners, especially above 0.6 Tm where high temperature creep occurs. Alloying ele-ment additions are bene�cial to increase strength until the solubility limit. Beyond this limit, another strengthening mechanism, precipitation strengthening, becomes effective (Brooks 1982, Heubner 1998).

1.3.1.2 Precipitation StrengtheningAs alloying elements are added to the matrix and solute atoms surpass their solubility limits, a new chemical constituent, a precipitate, will form. These chemically stable precipitates form the second phase in microstructures. In fact, a superalloy may contain a number of phases; some may even have more than eight different phases. The degree of precipitation hardening is closely related to the properties of the matrix and the new phases. Generally, to obtain optimized results, the matrix should be continuous, soft, and ductile while the precipitate phase should have appropriate distance, hardness, shape (morphology), and volume fraction. Controlling these parameters leads to control of both ductility and high strength.

Precipitates of high hardness obstruct the dislocation movement and therefore raise the required stress to plastic deformation. Dislocations can only continue their movement by cutting or climbing

FIGURE 1.4 Solid solution hardening mechanism.

TABLE 1.2Atomic Diameter Differences of Some Solid Solution Hardener in Nickel-Based Superalloys

Solute Element Co Fe Cr V Mo Zr Ta W

% ~Difference in atomic diameter (compare to nickel)

+1 +2 +0.5 +5 +9.6 +30 +15 +10

Source: Data from Askeland, D.R. 1985. The Science and Engineering of Materials. PWS Engineering: USA, p. 554.




on precipitates, which require much more stress. As such, hardness, yield, and tensile strength of the alloy are increased. If precipitated phases retain their stability and properties at high temperature, creep and stress rupture resistance of the alloy are also improved.

Unfortunately, not all phases that may form in superalloys always have positive contributions to properties. Depending on the number of alloying elements and service temperature, phase transfor-mations may occur and unwanted phases may precipitate and cause degradation in desired prop-erties such as creep and fatigue resistance. The main phases, which have important properties of superalloys, are listed below:

1.3.1.2.1 γ ′ PhasePrecipitation hardening is the main strengthening mechanism of the nickel and iron–nickel superal-loys. This hardening mechanism has a very limited effect on cobalt-based superalloys. Nickel and iron–nickel alloys mainly obtain their strength from the intermetallic phase γ ′ or Ni3(Al Ti). The γ ′ phase is very stable at high temperatures (up to at least 1100°C), ductile, and coherent with matrix (Sims et al. 1987). Ni–Al binary phase diagram with γ ′–Ni3Al is given in Figure 1.5.

The γ ′–Ni3Al phase also has an FCC (L12) structure like the γ matrix. As can be seen in Figure 1.6, nickel atoms are located on the faces of the ordered unit cell whereas aluminum atoms are located at the corners. Ordered structure means both aluminum and nickel atoms have �xed positions in the unit cell. Ordered structure is a very important property like coherency to increase the strength. The dislocation movement requires more energy in ordered structures than random structures (Reed 2006).

Although the γ ′–Ni3Al composition range is limited between 23% and 27% aluminum, both ele-ments may be substituted independently. The main substitute for aluminum is titanium with larger atomic radius. Niobium can substitute aluminum. Nickel can also be substituted with small atoms like cobalt and iron. The γ ′ yield stress increases as the temperature increases between −196°C and approximately 800°C. The effect of alloying element on yield strength of γ ′ can be seen in Figure 1.7. Temperature and stress level shifts in this �gure indicate that the γ ′ phase itself can also extensively bene�t from solid solution hardening.

The lattice parameter “a” of the pure nickel γ matrix is 0.3517 nm and γ ′–Ni3Al is 0.3570 nm, only about 1.5% larger than γ. Therefore, lattice strain between γ and coherent γ ′ (pure Ni3Al) is low. This coherency gives the opportunity for more precipitation. The aluminum/titanium ratio is an impor-tant parameter to control lattice strain hardening and is inversely proportional to lattice mismatch. High mismatch strain brings low-temperature strength and short-time high- temperature strength whereas low mismatch means high creep resistance due to stabler, slowly growing precipitates being able to form with high coherency. Therefore, the designers can select proper alloying accord-ing to service requirements and conditions. Lattice strain also affects the γ ′ morphology. In case of low strain (below 0.2%), γ ′ precipitates have spherical form; high lattice strain (about 0.5–1%) results in cubical γ ′ precipitates (Brooks 1982).

The γ ′ phase also has defects like γ matrix, which are point, line, and planar defects. One of the planar defects between γ and γ ′ caused by forced forbidden bonds between Ni–Ni and Al–Al is called antiphase boundary (APB). APBs form when the dislocation movement causes a slip in the γ ′ precipitate. APB has a dramatic effect on the dislocation movement in γ ′ and causes great increase in strength. Its effectiveness depends on the precipitate size, and the direction of the crys-tallographic plane therefore shows anisotropy (Reed 2006).

As the volume fraction of the γ ′ phase is increased, the strength of alloys is increased. This situation could be seen in Figure 1.8 for stress rupture strength at different temperatures as a function of γ ′ volume percent. The morphology of the γ ′ also transforms from spherical to cubical with increasing volume fraction of the γ ′ phase. Usually, at <25%, γ ′ is spherical, while at more than 35%, it becomes cubical. Early superalloys had low volume fraction of γ ′. Newly produced precipitation-strengthened superalloys contain more and more γ ′ for higher strength.



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Nickel, at.%2010

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ture

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ture

, °F

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200Al

Al

ε

β

γ

γ

δ

6.1 640°C

660.37°C

28 42 55

855°C

1135°C

1397°C, 83% 1387°C88.8%

1455°C

1640°C, 68.5%

5944 62

10 20 30 40 50Nickel (wt%)

60 70 80 90 Ni

Curietemperature

358°C

30 60 70 80

2910

1830

1470

1290

1110

930

750

570

390

86.7%

FIGURE 1.5 Nickel–aluminum phase diagram. (Reprinted with permission of Stoloff, N.S. 1990. ASM Handbook Vol. 1. Properties and Selection Irons Steels and High Performance Alloys: Wrought and P/M superalloys. ASM International: USA.)




The main limiting factor of this increase is the adverse effect of γ ′ on production techniques like weldability, forgeability, and so on. Superalloys with volume fraction of γ ′ of more than 40–45% (60% in powder metallurgy) are practically not forgeable and are generally produced by investment casting. Today, wrought nickel-based alloys contain 20–45 vol.% γ ′, while some cast alloys like Udimet 700 can contain up to 70 vol.% γ ′ and even some alloys 80 vol.% γ ′ (Durand-Charre 1997).

0 200 400 600 800 1000 12000

50

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150

2002190183014701110750

Temperature (°F)

Temperature (°C)

Flow

stre

ss, k

si39030

13806 at % Nb10.5 at % Ti + 2 at % Cr10.5 at % Ti2 at % CrNi3Al

1035

690

345

0

FIGURE 1.7 Temperature dependence and stress level shifts with in�uence of various solutes in γ with the permission of γ ′. (Reprinted with permission of Davis, F.R. 1997. ASM Specialty Handbook Heat Resistant Materials, p. 230, ASM International: USA.)

FIGURE 1.6 (See color insert.) Schematic view of the ordered L12-structured FCC cell.



The γ ′ particle diameter can be controlled via aging heat treatment. Through proper selection of γ ′ volume fraction and size, optimized strength can be obtained. Different heat treatment techniques (double or triple aging) are used sometimes to produce a combination of small- and large-size γ ′ pre-cipitates in the same alloy. Generally, as the distance between precipitated particles decreases, strength increases up to a certain distance. Further decrease in distance results in decrease in strength. Generally speaking, an alloy with low volume fraction of γ ′ requires one-step aging, while high volume frac-tion of γ ′ requires two or three steps. The γ ′ phase can also be formed as grain boundary �lm in some alloys. This morphology has a bene�cial effect on creep and stress rupture life. When γ ′ envelopes the grain boundary carbides such as Cr23C6, the ductility of the alloy greatly increases (Sims et al. 1987).

1.3.1.2.2 γ ″ PhaseAside from the γ ′ phase, some iron–nickel-based superalloys bene�t from ordered body centered tetragonal (BCT) D022-structured metastable γ ′′–Ni3(Nb,Ta) phase. The γ ′′ phase is able to with-stand high stress but remains stable at lower temperatures below approximately 650°C. Above this temperature, its strength sharply drops due to solutioning, coarsening, and transformation to the stable orthorhombic Ni3Nb-δ phase, which has a detrimental effect on strength. Orthorhombic Ni3Nb has acicular platelike morphology.

The γ ″ phase is coherent with the FCC matrix like γ ′ but it has a higher lattice mismatch (about 2.9%) and a disk-shaped morphology. Few iron–nickel alloys contain the γ ″ phase, Inconel 718 is the best representative, which acquires strength primarily from this phase. This alloy contains 5% niobium and 0.5% aluminum; therefore, both the γ ″ and γ ′ phases are present but γ ′′ is the primary strengthening agent with γ ′′-to-γ ′ ratio of 2.5–4. The aluminum amount is critical in alloys like Inconel 718 due to the solubility of aluminum in γ ′′.

Body-centered tetragonal lattice parameters of γ ′′ are a0: 0.3624 nm and c0: 0.7406 nm (Sims et al. 1987). Some modi�ed (Al + Ti)/Nb ratio alloys have a special form of cubic γ ′ particles coated with γ ′′ shell. This special morphology is obtained by careful heat treatment and shows a slow rate of coarsening.

0 15 30γ ′ volume %

45 60 750

29

58

87

116

1451000705°C

870°C980°C

760°C

800

600

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200

Stre

ss fo

r 100

-h li

fe (M

Pa)

Stre

ss fo

r 100

-h li

fe (k

si)

0

FIGURE 1.8 γ ′ volume percent versus stress rupture strength for nickel-based superalloys. (Reprinted with permission of Davis, F.R. 1997. ASM Specialty Handbook Heat Resistant Materials, p. 230, ASM International: USA.)




1.3.1.2.3 CarbidesSome elements in superalloys react with C and form hard carbides. Carbides form both in grain and on grain boundaries and serve as grain boundary strengthener and precipitation hardener. The effect of carbides on the properties of the �rst-generation superalloys was not fully understood during those years and C amount of alloys was kept at a minimum to avoid carbides. Although the bene�cial effect of carbides is better understood today, the morphology and location of the carbides in the microstructures should be carefully adjusted. Particularly, grain boundary carbides with thick continuous �lm morphology and intragranular carbides with widmanstatten morphology (a special type of acicular morphology) may decrease ductility. Although carbides are not as effective as γ ′ or γ ′′, they are the main precipitation hardener phase in cobalt-based superalloys.

The main types of carbide seen in superalloys are MC, M23C6, M6C, and less frequently M7C3. MC carbides form �rst, with the reaction between reactive elements such as Ti, Hf, Ta, Nb, and C during solidi�cation. Other types of carbides mainly form from the transformation of MC carbides during service or heat treatment. Other types of carbides may also form by a combination of γ ele-ments and C. MC carbides have FCC structure and locate in random positions like interdendritic, intergranular, or transgranular sites. Hf and Nb carbides are very stable. The basic reactions of MC carbides are given below with an example reaction.

M + C → MC (1.1)

(Ti, Hf, Nb, W, Ta, Mo) + C → (Ti, Hf, Nb, W, Ta, Mo)C (1.2)

The optical micrograph of TiC can be seen in Figure 1.9 with cubic morphology and its char-acteristic yellowish color in Hastelloy X alloy. M23C6 carbides are formed by the decomposition of MC and M6C carbides and the reaction between solute C in γ and former carbide elements. The transformation reactions of MC and M6C carbides to M23C6 carbides are given with examples below.

MC + γ → M23C6 + γ ′ (1.3)

(Ti, Mo)C + (Ni, Cr, Al, Ti) → Cr21Mo2C6 + Ni3(Al, Ti) (1.4)

M6C + M′ → M23C6 + M′′ (1.5)

Mo3(Ni, Co)3C + Cr ↔ Cr21Mo2C6 + (Ni, Co, Mo) (1.6)

FIGURE 1.9 Cubical TiC carbide microstructure in wrought nickel-based solid solution superalloy Hastelloy X (magni�cation 500×).



As can be seen in the reactions, the dominant element in Mr23C6-type carbide is Cr, Mo, W, Co, and Fe may also form with limited amount in this carbide structure. Because superalloys generally contain Cr, M23C6 is frequently observed in superalloys. M23C6 carbide is almost always found in grain boundaries but sometimes they may also be found in twinning bands, stacking faults, which are suitable precipitation areas with their relatively low orders. Grain boundary carbides like M23C6 increase stress and creep rupture resistance of superalloys by impeding grain sliding and retarding cavity formation. This carbide group also makes possible the stress relaxation of alloys under high temperature via allowing freer expansion of grains. Grain boundary carbides morphology dramati-cally affects the alloy properties. One of the ideal morphologies for grain boundary carbides is thin, zipper-like shape, as shown in Figure 1.10. This morphology is very effective in preventing grain boundary sliding at high temperatures (Sims et al. 1987, Stoloff 1990).

Grain boundary carbides grow with time and temperature and form excessively thick, continu-ous grain boundary carbide network, which causes brittleness and denuded zones around grain boundaries. The basic mechanism of formation of denuded zone is increasing the solubility of γ matrix due to their decreasing Cr content. This situation causes γ ′ elements solutioning in γ, hence decreases in strength and volume fraction of γ ′ near grain boundaries.

The γ ′ phase also forms in reaction 1.3 and 1.4. The γ ′ phase envelops the M23C6 carbides as a �lm through these reactions and thus slows down or stops the unwanted growth of M23C6 carbides. Therefore, both unwanted growth of M23C6 carbides and the creation of denuded zones is hindered. M23C6 carbides have a complex cubic structure which is very close to the TCP (topologically closed phase) structure and unwanted σ plates usually form on these carbides.

M6C carbides may form in superalloys containing more than 6% Mo plus W. The chemical com-position may vary but Mo3(Ni, Co)3C and (Ni, Co)2W4C are frequently observed. M6C carbides also have complex cubic structure like M23C6. While the M23C6 structure is similar to the σ phase, the M6C phase is similar to another TCP phase µ. The increasing amount of Mo and W in superalloys may cause undesirable formation of µ phase.

Due to its stabler structure when compared to M23C6, M6C is a very useful tool to control grain size in wrought alloys during plastic working. Besides 1.5 and 1.6, reactions that cause M6C are given below as 1.7 and 1.8.

MC + γ → M6C + γ ′ (1.7)

(Ti, Mo)C + (Ni, Co, Al, Ti) → Mo3(Ni, Co)3C + Ni3(Al, Ti) (1.8)

M7C3 carbides are rarely observed types and may be found especially in Co-based superalloys. It precipitates on grain boundaries with blocky morphology and thus may prevent the formation of M23C6 at grain boundaries.

FIGURE 1.10 Grain boundary zipper-like M23C6 carbides microstructure in Rene 41 wrought nickel-based superalloy (magni�cation 1500×).




Another important bene�t of carbides is stopping or slowing down the formation of some detri-mental phases as the reactive elements are bound with C to form carbides (Stoloff 1990).

1.3.1.2.4 η PhaseEta phase (Ni3Ti) is found in superalloys with high titanium/aluminum ratios when alloys are exposed to high temperature under long service time. Its structure is hexagonal close packed and has a detrimental effect on alloy properties especially notched stress rupture strength and creep ductil-ity. η phase can be formed on both grain boundary with γ + η lamellar structure and intragranular sites with needlelike platelets.

1.3.1.2.5 σ Phaseσ phase is often observed in iron–nickel- and cobalt-based alloys. Its chemical composition is quite wide but crystal structure is only tetragonal. Some forms of this phase play a detrimental role while others only show a minor effect or are effectless on properties. When it precipitates on globular intragranular particles, it even improves creep resistance. The worst-affected property is high-tem-perature stress rupture life. It is believed that σ phase causes alloy depletion damage of the γ matrix. FeCr, FeCrMo, and CrNiMo are some of the chemical compositions of σ. A more general formula of σ is (Fe,Mo)x(Ni,Co)y. The σ phase is observed when superalloys are subjected to long service time between 540°C and 980°C. When the σ phase becomes platelike, it would serve as crack initiation site and propagation path, causing low-temperature brittleness.

1.3.1.2.6 Laves PhaseLaves phase is of TCP structure and most commonly found in iron–nickel-based superalloys. The chemical composition is AB2 where A represents Fe and less commonly Cr, Mn, and Si, and B rep-resents Mo, Ti, and Nb in iron–nickel alloys. Some of the cobalt-based alloys also have Laves phase with a chemical composition of Co2W, Co2Ta, and Co2Ti. Laves phase has a detrimental effect on ductility when the precipitates are too much. Like many other phases in superalloys, Laves phase formation is also promoted by high temperature, time, and stress.

Generally speaking, TCP phases like σ and Laves are undesired because they cause brittle-ness, reduce creep, and stress rupture life. Special efforts have been made about these phases to determine forming conditions (time, temperature, stress, chemical composition, etc.) and effect on properties. One method is using computer simulation to predict the possibility of the formation of unwanted TCP phase according to types of elements and their amount in the alloy. PHACOMP (phase computations) may be the best known program, which calculates the net Nv—the number of electron vacancies, that is, the number of unpaired d shell electron of each alloying element in a given alloy. This program also takes into account the amount of elements in the alloy. The Nv value is calculated for residual γ matrix after the formation of possible precipitates like carbides, γ ′, and so on. As the Nv value for residual γ matrix exceeds about 2, �ve TCP phases can form. Other programs like “Sigma Safe” based on phase equilibrium are also used by alloy designers. BCC-structured alloying elements like Nb, Cr, W, and Ta cause the formation of TCP phases. Therefore, alloys con-taining a high amount of these elements tend to form TCP phase (Sims et al. 1987).

1.3.1.3 Effect of Grain-Related Properties on StrengtheningOther important factors affecting the strength of superalloys are the grain size and the ratio of grain size to component thickness. Generally, superalloys that have coarse grains have better creep and stress rupture resistance. But coarse grains also mean lower tensile strength. The proper grain size for a particular component may also change under service conditions. At low to intermediate tem-peratures, the in-grain tensile strength is lower than grain boundaries. As the temperature increases, the in-grain strength �rst reaches (equi-cohesive temperature) and then exceeds the grain boundary strength. As grain size increases, the total amount of grain boundary reduces. Thus, the desired amount of grain boundaries can be obtained by adjusting grain size with proper heat treatment.



Grain boundary-related failures in the superalloys can be reduced by con�guring grain orientation to minimize transverse grain boundaries in unwanted direction. Grain orientation can be obtained by different methods such as rolling for wrought alloys and directional solidi�cation for cast alloys. The strength of the turbine blades and vanes is signi�cantly improved this way. The crystallographic ori-entation also considerably alters strength, desired <001> orientation, which has maximum strength in FCC lattice, can be set parallel to the main stress axis in turbine blades by directional solidi�cation (Reed 2006). Another extreme way of reducing grain boundary-related failures is completely elimi-nating them with single crystal technology. Single crystal superalloys can be produced by application of very sensitive process parameters. These alloys have considerably high creep and thermal fatigue resistance. The main application of single crystal superalloys is turbine blades and vanes. There is still a huge production of polycrystalline (PC) superalloys in various parts.

1.3.1.4 Effect of Alloy ChemistryThe effect of alloying elements on various superalloy properties was brie�y explained in the previ-ous sections. In this section, other elements that were not mentioned before are explained. High-melting-point refractory elements such as B and Zr are widely used in superalloys to control grain size and increase grain boundary strength. The selective oxidation of Cr and Al elements in the alloys creates protective Cr2O3 and Al2O3 oxide layers under hot service environment, preventing further oxidation into the interior of the alloy. To be effective, ideally, the protective oxide layer should have the following properties:

• Stable service environment• Slow growth into the alloy• Dense, thus able to block direct oxygen passage• Resistant to thermal and mechanical stress (Mevrel and Veyes 1989)

The use of Cr2O3 scale is limited at low temperature as it changes to volatile oxide above 900°C, especially under high-velocity gas stream. Therefore, for protection at temperatures higher than 900°C, Al2O3 scale is still the only choice.

Another critical factor that causes the reduction of oxidation resistance and mechanical strength is the impurity in the alloy chemistry. The impurities form unwanted phases in the system, thus causing property deterioration. The reliability is signi�cantly enhanced by virtue of advances in vacuum metallurgy and melting practices which helped the elimination of trace elements such as S, P, Si, O, N, Pb, Se, and so on. These detrimental elements usually cause brittleness and decreas-ing incipient melting point of the alloy. Till now, there is no information on the exact maximum permissible amount of each trace element in each superalloy, as these amounts vary from alloy to alloy (Heubner 1998). On the other hand, however, some trace elements are intentionally added for a speci�c purpose. Care should be taken in this case so as not to go over the limit. For example, boron and zirconium are believed to �ll the vacancies and lattice imperfections at or near the grain boundaries, thus inhibiting the growth of grain boundary M23C6 carbides for signi�cant improve-ment of stress rupture and creep resistance (Sims et al. 1987). However, B addition is favorable up to 50 ppm; beyond this limit, the formation of low-melting point borides causes embrittlement of grain boundaries (Heubner 1998).

Table 1.3 provides a summary of chemical composition effect on superalloys.

1.3.2 IRON–NICKEL-BASED SUPERALLOYS

Iron–nickel-based superalloys are chronologically the �rst superalloys group and were developed from austenitic stainless steels. The �rst alloys were produced via the addition of titanium to high-chromium-content austenitic stainless steels to obtain age hardening. German Tinidur and the very popular American A-286 alloy is an example. These alloys contain a high amount of iron and their




amount of precipitates is quite low. They required a minimum of 25% nickel to stabilize FCC aus-tenitic matrix. These alloys contain 25–60% nickel and 15–60% iron (Brick et al. 1977). Iron is a low-cost element and is found abundantly. Therefore, special efforts have been made to improve this group of superalloys. Although high iron content provides considerable cost saving and improves alloy forgeability, oxidation and corrosion resistance decrease its signi�cance.

Iron-based superalloys have similar strengthening mechanisms to nickel-based superalloys. Mo and Cr are used to obtain solid solution hardening. Rarely, W is also used as a substitute for Mo with the disadvantage of high cost and density. Besides solid solution hardening, Cr also provides effec-tive oxidation and sul�dation resistance in the service temperature range of this group of alloys. The amount of Cr should be suf�ciently high to form protective continuous oxide layer on the alloy sur-faces for effectiveness. Studies show that Cr amount should be higher than 9% wt. This value is well below the Cr content in most iron–nickel-based superalloys. Carbon, nitrogen, and boron elements with small atomic radius also improve solid solution hardening via interstitial placement in FCC lat-tice. As the solid solution hardener elements also cause decrease in solubility of the FCC matrix, the required amount of precipitation hardener elements are thus signi�cantly decreased. Ti is added to the matrix to form coherent γ ′ precipitation hardener phase as in nickel-based superalloys. However, Al addition is very limited compared to nickel-based superalloys due to the electronic structure of iron, which causes the formation of metastable structure of γ ′ and then transforms to useless or other detrimental stable phases under working conditions (Durand-Charre 1997). As such, Al is usually used in iron-based superalloys only as a deoxidizing agent. The lattice parameter of the FCC matrix increases with iron content. Using Ti instead of Al causes a similar effect on γ ′ phase. Therefore,

TABLE 1.3Role of Elements on Superalloys

Effect Iron-Based Cobalt-Based Nickel-Based

Solid solution strengtheners Cr, Mo Nb, Cr, Mo, Ni, W, Ta Co, Cr, Fe, Mo, W, Ta

FCC matrix stabilizers C, W, Ni Ni . . .

Carbide formation:

MC Ti Ti, Ta, Nb W, Ta, Ti, Mo, Nb

M7C3 … Cr Cr

M23C6 Cr Cr Cr, Mo, W

M6C Mo Mo, W Mo, W

Carbonitrides

M(CN) type C, N C, N C, N

Forms γ ′ Ni3 (Al, Ti) Al, Ni, Ti … Al, Ti

Retards formation of hexagonal η(Ni3Ti) Al, Zr … …

Raises solvus temperature of γ ′ … … Co

Hardening precipitates and/or intermetallics Al, Ti, Nb Al, Mo, Tia W, Ta Al, Ti, Nb

Forms γ ″ (Ni3Nb) … … Nb

Oxidation resistance Cr Al, Cr Al, Cr

Improves hot corrosion resistance La, Y La, Y, Th La, Th

Sul�dation resistance Cr Cr Cr

Increases rupture ductility B B, Zr Bb, Zr

Causes grain-boundary segregation … … B, C, Zr

Source: Stoloff, N.S. 1990. ASM Handbook Vol. 1. Properties and Selection Irons Steels and High Performance Alloys: Wrought and P/M Superalloys. ASM International: USA, p. 2305.

a Hardening by precipitation of Ni3Ti also occurs if suf�cient Ni is present.b If present in large amounts, borides are formed.



high Ti amount leads to a decrease in γ/γ ′ lattice mis�t. Ni serves as an austenite stabilizer and γ ′ former (Ni3Ti–Al) in iron–nickel-based superalloys. As mentioned before, this alloy group also ben-e�ts from other precipitating phase, γ ′′ (Ni3Nb). This phase has a BCT structure which signi�cantly improves the mechanical properties of the alloy. The main drawback of γ ′′ is that, when compared to γ ′, its useful upper temperature limit is considerably lower, only about 650°C. Therefore, the iron–nickel-based superalloys are good in high-strength, intermediate-temperature regime.

Grain boundary carbides, such as Cr23C6, improve grain boundary strength and IGO resistance. It is determined that the unwanted MC carbides of thick continuous morphology form on grain boundaries and cause brittleness in alloys such as A-286 and Incoloy 901. Other important carbide-forming elements are Ti (in MC form) and Mo (in M6C form) in iron–nickel-based superalloys. B and Zr are also frequently added to the iron–nickel-based alloys to improve stress rupture and hot workability.

Iron–nickel-based superalloys can usually be classi�ed into four subgroups. These are γ ′-strengthened alloys (like nickel-based superalloys), primarily γ ′′-strengthened alloys, solid solu-tion alloys, and low-expansion alloys.

γ ′-Strengthened alloys evolve from austenitic stainless steels with the addition of Ti. Therefore, early examples of this group contain an iron-rich matrix whereas later alloys contain more nickel than iron and are called nickel-rich subgroup. The well-known examples of iron-rich group are A-286, Tinidur, Discaloy, and V-57. The chemical composition of V-57 is very similar to A-286 with only slight adjustments and sometimes referred to as super A-286. These alloys contain less solid solution and lower volume fractions of precipitate strengtheners than nickel-rich alloys. Their service temperature range is limited to about 650°C. Nickel-rich iron–nickel-based superalloys can withstand higher temperature and stress levels than their iron-rich counterpart. Also, they have relatively low cost advantage. Inconel X-750 and alloy 901 are typical examples of this subgroup.

γ ′′-Strengthened alloys are widely used in the gas turbine industry. Inconel 718 is a member of this group; it was invented in the 1950s and perhaps is the most popular alloy with a production rate roughly half of all superalloys in the world. These nickel-rich group alloys contain both γ ′ and γ ′′ precipitates. Because of its limited temperature range but perfect stress characteristic, Inconel 718 is so widely used in gas turbines that, for instance, nearly one-third of the materials of GE CF6 engine consist of Inconel 718. This alloy is utilized mainly at gas turbine parts where high stress and a relatively low-temperature environment are present such as rear stages of compressor blades and vanes, shafts, supports, cases, and so on.

Solid solution alloys contain little or no precipitates and in turn their strength is quite low but they have very good oxidation resistance. 19-9 DL, Alloy N-155 (Multimet) are well-known exam-ples of solid solution alloys. Because of their low strength and high oxidation resistance properties, these iron-rich alloys are used in highly oxidizing but low-stress environment such as combustion chambers and �ame holders in gas turbines.

Low-expansion iron–nickel-based alloys are used in gas turbines where tight clearance control is required especially between rotating and static parts. Incoloy 903 and Pyromet CTX-1 are typical examples. These iron-rich alloys contain Co and also have very good strength characteristic owing to γ ′ and less frequently η (Ni3Ti) precipitates but are able to retain their strength up to 650°C. Low expansion characteristics stem from the absence of ferrite stabilizers like Cr and Mo. This group suffers from poor oxidation resistance due to lack of Cr; thus, it easily gets oxidized beyond 550°C. The main approach to overcome the oxidation problem is to apply oxidation-resistant coating on these alloys.

1.3.3 NICKEL-BASED SUPERALLOYS

Nickel-based alloys are the strongest group of all superalloys in the world in both strength and temperature capability; therefore, their usage has vital importance in designing high-performance gas turbines. These alloys are used extensively in gas turbines where most demanding applications




are needed such as high- and low-pressure turbine blades, vanes, disks, and so on. The key fac-tor of nickel-based superalloys’ strength can be explained with precipitation hardening. It is the most useful mechanism at high temperatures and nickel-based alloys bene�t from this mechanism most effectively through γ ′ phase. Although iron–nickel-based alloys also contain γ ′, their service is limited at intermediate temperatures as explained in iron–nickel-based superalloys. Comparing wrought and cast nickel-based superalloys, the cast one has better high-temperature strength such as stress rupture and creep life.

The evolution of nickel-based superalloys goes back to solid solution strengthening. The Ni–Cr phase diagram is shown in Figure 1.11. From the phase diagram, Cr is quite soluble (around 35%) in Ni. Nichrome alloys are generally used as electrical resistance wire. One of the well-known alloys of this class is Nichrome V with 80–20% Ni–Cr chemical composition that provides a proper bal-ance between strength, oxidation resistance, and ductility. Aluminum and titanium are added to this class of alloys to obtain additional strength and corrosion resistance. Nimonic 80 alloy was developed in England in 1941 and is one of the �rst precipitation-hardened nickel-based superalloy with the addition of 2.25% titanium and 1% aluminum to Nichrome V with 20% chromium ratio remaining �xed (Smith 1993). Many alloying elements have been added to nickel-based superalloys with time to improve different properties. Also, a lot of solid solution nickel-based superalloys exist that evolved from Ni–Cr alloys such as Inconel series 600, 601, 625.

Nickel-based superalloys can be subdivided into three categories: solid solution, precipitation-hardened, and oxide dispersion-strengthened alloys.

Austenitic solid solution alloy Inconel 600 (Ni–15.5%Cr–8%Fe) shows very good corrosion and oxidation resistance. This alloy cannot be hardened with heat treatment but could be hardened by cold work. Inconel 601, Ni–23%Cr–14.1%Fe–1.4%Al, has better corrosion and oxidation resistance than Inconel 600 with higher Cr amount. Aluminum adds extra oxidation resistance via the for-mation of a protective alumina layer. Both alloys are widely used in heat treatment furnaces, heat exchangers, and so on. Hastelloy X is also in this group and is very popular. Hastelloy X can be

5000 10 20 30 40 50 60 70 80 90 100

0 10 20 30 40Atomic percent chromium

50 60 70 80 90 1001663°C

CrWeight percent chromiumNi

(Ni)

γ590°C

(Cr)

L

1345°C

700

900

1100

1300

Tem

pera

ture

(°C

)

15001455°C

1700

1900

FIGURE 1.11 Ni–Cr phase diagram. (Reprinted with permission of Nash, P. 1992. ASM Handbook Vol. 3 Alloy Phase Diagrams, p. 692, ASM International: USA.)



easily welded and thus is used extensively in gas turbine and combustion chamber up to 1200°C. Similar to its iron–nickel-based solid solution counterparts, nickel-based solid solution superalloys lack high strength but possess good oxidation and corrosion resistance.

Early examples of precipitation-hardened alloys like Nimonic 80 had a low amount of γ ′ pre-cipitate, while newer ones have increased aluminum and titanium concentration to increase γ ′. As more alloying elements are introduced, some side effects have been observed, such as the formation of unwanted phases. Alloy designers solve these problems by using new techniques such as special heat treatment, and sometimes even addition of more alloying elements.

Precipitation-hardened nickel-based alloys are the strongest superalloys and extensive effort has been made for years to obtain even higher strength. Many important milestone improvements like directional solidi�cation and single crystal technology are the result of these efforts (to be described in detail later in production techniques section). DS was used to produce SM-200 (will turn into MAR-M 200 later) nickel-based cast alloy turbine airfoil in 1960s. Today, other nickel-based cast superalloys like Rene 80 used extensively in hottest turbine section parts such as high-pressure tur-bine blade are produced with directionally solidi�ed technique. Rene 80 contains Al and Ti for γ ′ formation, Mo and W for γ solid solution strengthening, and C, Zr, and B for grain boundary re�ne-ment and ductility. This alloy shows excellent stress rupture combined with elevated temperature ductility up to 1000–1050°C. Its hot corrosion resistance is also very good with 14% Cr and high Ti/Al ratio (5/3). It has a very stable microstructure in long-term applications owing to low vacancy. Although this alloy has good resistance to oxidation and hot corrosion, in gas turbines it generally still needs coating to protect from service atmosphere. Rene 80 roughly contains 47% vol. γ ′, 2% vol. MC, with minor amount of M23C6 and M3B2 after multistep aging and full heat treatment. MC carbides are rich in Ti, W, and Mo and then turn into M23C6 and M6C under service conditions. The microstructure of Rene 80 in fully aged alloys can be seen in Figure 1.12. Hafnium is added to both MAR-M 200 and Rene 80 to improve grain boundary ductility. “Rene 80 Hf” contains about 0.75% Hf with slightly reduced Ti, C, and Zr. Since the appearance of MAR-M 200 in the late 1960s, Hf is a common element for cast alloys and is found to be very bene�cial to improve key grain boundary properties like creep, stress rupture (Sims et al. 1987).

Udimet 700 is another important nickel-based precipitation-hardened alloy as wrought and cast. Forging is quite dif�cult and must be performed under optimum temperature ranges to avoid surface cracking. Machining is also dif�cult but possible under all heat treatment conditions. Fully aged con-dition gives the best results with carbides. It has outstanding stress rupture, creep, oxidation, and hot corrosion resistance at elevated temperatures. It forms continued and adherent oxide layer because

FIGURE 1.12 Microstructure of cast Rene 80 nickel-based superalloy after fully aged heat treatment. Script pattern carbide morphology and shrinkage cavities, occurred during casting can be seen (see arrows) (magni�cation 100×).




of high amounts of Cr and Al which protect the alloy from further oxidation at 1000–1050°C under continuous or at 925°C under intermittent service conditions. A high amount of Al + Ti (4.5 + 3.5%) in this alloy allows a high amount of γ ′ after multistep aging treatment. It is used widely as gas turbine hot section parts such as turbine buckets where both stress and high-temperature resistance are needed. Wrought alloy Astroloy has the same basic composition with slightly reduced Co, Al, and B. Both Udimet 700 and Astroloy tend to undergo sigma phase formations during long service at elevated temperatures. Rene 77 is an answer to this problem with carefully designed chemical compositions to result in low electron vacancy number per atom. 1000 h rupture strengths of some wrought nickel-based superalloys, including Udimet 700 and Astroloy, are given in Figure 1.13. It is noteworthy that solid solution alloys like Hastelloy X and Inconel 601 are also given in this �gure and their low strength characteristic can be easily distinguished.

7000

50

100

150

200

250

300

350

400

Stre

ss (M

Pa)

Stre

ss (k

si)

450

500

550

600

650

7001400 1500 1600 1700 1800

Temperature (°F)1900 2000 2100 2200

100

90

80

AstroloyD-979Hastelloy XIN 100 gatorizedInconel 600Inconel 601Inconel 625Inconel 718Inconel X-750M-252MA 754MA 6000Nimonic 80A

Udimet 500Udimet 520Udimet 700Udimet 710Udimet 720Udimet AF2-IDAWaspaloy

René 41

70

60

50

40

30

20

10

0800 900 1000

Temperature (°C)1100 1200

FIGURE 1.13 1000 h rupture strengths of selected wrought nickel-based superalloys. (Reprinted with permis-sion of Davis, F.R. 1997. ASM Specialty Handbook Heat Resistant Materials, p. 234, ASM International: USA.)



Although solid solution alloys can be more easily produced in both cast and wrought forms, precipitation-hardened alloys with high volume fraction of γ ′ can usually only be produced in cast-ing or powder metallurgy route owing to their poor fabrication characteristics such as welding, forging, and so on.

Although the idea is old, ODS alloys could practically be produced only after the development of mechanical alloying process in the 1970s (Tien and Caul�eld 1989, Wilde 1975). Since then, powder metallurgy is applied to superalloys with high volume fraction of ODS unsuitable for conventional forging. The starting material comes from ingot metallurgy (Campbell 2006), and thus suffers some technical limitations; other than that, powder metallurgy has a clear advantage on material saving.

With increasing temperature, the microstructural compounds of superalloys, such as precipita-tion hardener γ ′′ phase, that has the highest impact on mechanical strength, would lose their stabil-ity and become part of a solid solution or transform to other stabler phases. Some early efforts of addition of Yttria (Y2O3) in superalloys were not successful in obtaining �nely dispersed particles (Tien and Caul�eld 1989).

Mechanical alloying (MA) is based on continued fracturing and rewelding of powder particles in a ball mill until desired �ne distributions of oxide particles are obtained. These �nely dispersed, hard, stable particles serve as a strengthening agent by obstructing the dislocation movement.

MA 754 was the �rst commercial alloy. Similar to Nimonic 80, MA 754 evolved from 80-20 Ni-Cr Nichrome V alloy with the addition of about 1 vol.% Yttria and a small amount of Al, Ti, and C. Its microstructure consists of �nely dispersed yttrium aluminates and coarser carbonitrides in the γ matrix. The oxide dispersoids are always noncoherent with the matrix; therefore unlike γ ′ and other coherent precipitates, coherency-related strengthening mechanisms are not applicable to dispersoids. MA 754 is widely used in gas turbine nozzles, vanes, and other high-temperature applications. As other ODS alloys, MA 754 oxidation and hot corrosion resistance properties are very good because of high Cr content (Davis 1997). MA 758 was developed to obtain further cor-rosion resistance with even higher Cr content (30 vol.%). MA 758 is used in glass production. MA 6000 combines the advantages of strength from both γ ′ and oxide dispersoids. Al, Ti, and Ta form a high volume fraction of γ ′ (45–50%) and is responsible for intermediate temperature strength, whereas Yttria dispersions provide strength at even higher temperatures. Aluminum also signi�-cantly improves oxidation resistance together with chromium. W and Mo elements provide solid solution strengthening. MA 6000 is a perfect alloy for high temperature strength applications such as turbine blades and vanes.

1.3.4 COBALT-BASED ALLOYS

Cobalt-based alloys differ from iron–nickel- or nickel-based superalloys by the absence of ordered coherent precipitates γ ′ or γ ″. Although a few alloys such as CM-7 (Modi�ed L-605) contain Co3(Al,Ti), their useful temperature is very limited and not stable for long service. Cobalt-based superalloys extensively bene�t from carbides and solid solution strengthening to overcome this absence. Therefore, cobalt alloys contain much more C than other groups of superalloys. The �nely distributed carbides signi�cantly improve strength and these carbides are generally stabler than carbides in other groups of superalloys. Similar to nickel-based alloys, common carbide types in cobalt-based superalloys are M6C, M7C3, and M23C6. As cobalt-based superalloys do not contain Ta, Ti, Zr, or Hf (M in MC), MC-type carbide is not found in cobalt-based superal-loys. As the main elements for M6C are W and Mo, this carbide can form if suf�cient amount (5% at. or more) are present. M23C6-type carbides mainly form from Cr as in other superalloys and can also contain smaller quantities of W and Mo. M6C carbides may transform to M23C6 and precipitate on grain boundaries mainly as Cr23C6. These grain boundary Cr23C6 prevent the alloy from grain boundary sliding which in turn prolonges creep rupture life at elevated temperature (Smith 1993). Cr23C6 carbides also precipitate in stacking fault areas in the lattice and increase




strength by pinning dislocations. Unfortunately, embrittlement is the most important unwanted side effect of carbides in cobalt-based alloys that contain more carbides than other superalloys. This is especially pronounced when cobalt-based alloys are subjected to high-temperature expo-sure for a long time (Smith 1993).

Although the mechanical strength level is less than nickel-based superalloys, cobalt alloys have many distinctive properties over other groups of superalloys. Generally, cobalt-based alloys have better or comparable hot corrosion resistance due to higher Cr content, good creep, and stress rup-ture resistance at elevated temperature, thermal fatigue resistance, and weldability. Cobalt-based superalloys could also be melted in air at low cost as many cobalt alloys do not contain reactive elements such as aluminum or titanium.

Similar to iron–nickel-based superalloys, cobalt alloys also need suf�cient but lesser amount of minimum 10% of nickel to stabilize their austenitic FCC structure except some cast alloys such as Haynes 21, MAR-M 302. Fe, Mn, and C are also austenite stabilizers while W and Cr act as HCP stabilizers. Cr is the main alloying element, 20–30% in cobalt-based superalloys and improves oxi-dation resistance. Iron and nickel improve workability. Their microstructures are relatively simple, and the main solid solution strengtheners are Nb, Ta, W, and Mo which have large atomic diameter difference with cobalt matrix. MAR-M 302 cast alloy is a well-known and good example of solid solution-hardened alloy with a 10% W and 9% Ta. This alloy also bene�ts from carbide strengthen-ing and Zr and B additions to improve grain boundary properties, and is used in gas turbine parts at elevated temperature up to 1100°C. Alloying element additions to cobalt alloys also bring up unwanted TCP phases such as sigma and Laves. Cobalt-based alloys are more sensitive than nickel-based alloys in the formation of these unwanted phases. One of the examples that is subject to Laves (Co2W) phase formation happens in the very popular L-605 (also known as Haynes 25) cobalt-based alloy, in both cast and wrought forms. Silicon content with the presence of tungsten made suscep-tible L-605 to Laves formation under long life service conditions at elevated temperatures. Laves phase causes embrittlement in the alloys. Haynes 188 is an answer to this problem with controlled amount of silicon, reduced tungsten, and other balanced elements. The microstructure of cobalt-based Haynes 188 superalloy is given in Figure 1.14.

L-605 has good machinability, workability, and weldability. It is hardened by cold work and is widely used in industrial and aircraft gas turbine hot section parts and other areas such as combus-tion chamber, �ame holder components, liners, components of nuclear reactor, surgical implants, and so on (Davis 1997).

FIGURE 1.14 Microstructure of Haynes 188 cobalt-based superalloy. Arrows show twinning bands, which indicate plastic deformation during processing.



Cobalt-based alloys may be classi�ed according to their intended application area: high temperature, wear resistance, and corrosion resistance. All of these are found in both cast and wrought forms. Some corrosion-resistant alloys are also produced by the powder metallurgy route (Davis 1997).

L-605 and Haynes 188 are good examples of heat-resistant superalloys. Their �at stress rupture curves mean long time capacity of low stress at high temperatures. Stellite series 6, 21, 31 (also known as X-40) are well-known wear-resistant alloys. Stellite 6 is the most ductile and strongest among them. It is used not only for wear-resistant, but also for heat-resistant applications due to its capability of retaining hardness up to 1050°C. It is naturally hard with high C and is resistant against oxidation with high Cr content, and is thus used as wear pad in gas turbines and erosion shield in steam turbines. Some of the important properties of superalloys are given in Tables 1.4 through 1.7.

1.4 PRODUCTION AND PROCESSING TECHNIQUES FOR SUPERALLOYS

The evolution of superalloys depends not only on the development of new alloys and a deeper understanding of their microstructure, but also on the progression of their production and pro-cessing techniques. For metallic materials including superalloys, melting is the �rst step no matter what follows next: casting, forging, or powder metallurgy. Technologies of melting super-alloys have signi�cantly improved in the past decades. These improvements made possible more reliable, cleaner, and more precise control of chemical composition and better reproducibility. Today’s alloy designers feel much freer than in the past at utilizing the alloying elements they like.

Obviously, investment casting is the basic and primary method for the production of superal-loys that are hard to machine with near net shape and very low surface roughness. It is possible to obtain parts of complex internal cooling geometries with the advancements of investment casting that would otherwise be very dif�cult or impossible. DS and SX alloy production techniques have opened new horizons in turbine design.

Forging technology has gained new dimensions with isothermal forging and thermomechani-cal processing concepts. Application of superplastic forming technology to superalloys is also a signi�cant milestone in the development of forging techniques. As in many other superal-loys production, extensive use of computer technology drastically reduces cost and time in die design and determination of parameters. Computer technology has also greatly increased forged part quality thanks to precise material �ow characteristics achieved through �nite ele-ment simulations.

Although the fundamentals of powder metallurgy techniques have been known for a long time, their application to the superalloys is relatively late. The increasing demand of better strength in gas turbine engines meant more precipitation hardener in superalloys. The answer is vacuum invest-ment casting for turbine blades and powder metallurgy for turbine disks. Powder metallurgy makes the production of massive parts that contain high amount of γ ′ phase possible but at the same time makes forging dif�cult. Another interesting bene�t of powder metallurgy techniques is the devel-opment of a new superalloy group. With combined usage of powder production and mechanical alloying techniques, ODS superalloys that have a chemical composition that otherwise cannot be obtained have been obtained.

Regardless of the production route, superalloys usually need a combination of different fabrica-tion techniques to convert into complex turbine components. The components may involve cooling or fastening holes, may have surface to be machined to better roughness or joining sections which need welding or brazing, and so on. Furthermore, parts, which come from different production routes such as casting, forging, and so on, are joined to obtain the components. It is of critical value that these fabrication techniques must not have a physically or chemically detrimental effect on the component above an acceptable level.




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.

Nim

onic

115

6014

.313

.2..

...

...

.4.

93.

7..

...

...

.0.

150.

160.

04..

.

Nim

onic

263

5120

205.

9..

...

.0.

52.

1..

.0.

40.

30.

060.

001

0.02

...

Nim

onic

942

aba

l12

.5...

6..

...

.0.

63.

737

0.2

0.3

0.03

0.01

...

...

Nim

onic

PE

.11a

bal

18..

.5.

2..

...

.0.

82.

335

0.2

0.3

0.05

0.03

0.2

...

cont

inue

d



TAB

LE 1

.4

(con

tinu

ed)

Che

mic

al C

ompo

siti

on o

f Sel

ecte

d W

roug

ht N

i-B

ased

Sup

eral

loys

Allo

y

Com

posi

tion

(%

)

Ni

Cr

Co

Mo

WN

bA

lTi

FeM

nSi

CB

Zr

Oth

er

Nim

onic

PE

.16

4316

.51

1.1

...

...

1.2

1.2

330.

10.

10.

050.

02..

...

.

Nim

onic

PK

.33

5618

.514

7..

...

.2

20.

30.

10.

10.

050.

03..

...

.

Pyro

met

860

4312

.64

6..

...

.1.

253

300.

050.

050.

050.

01..

...

.

Ren

é 41

5519

111

...

...

1.5

3.1

...

...

...

0.09

0.00

5..

...

.

Ren

é 95

6114

83.

53.

53.

53.

52.

5..

...

...

.0.

150.

010.

05..

.

Udi

met

400

aba

l17

.514

4..

.0.

51.

52.

5..

...

...

.0.

060.

008

0.06

...

Udi

met

500

5418

18.5

4..

...

.2.

92.

9..

...

...

.0.

080.

006

0.05

...

Udi

met

520

5719

126

1..

.2

3..

...

...

.0.

050.

005

...

...

Udi

met

630

aba

l18

...

33

6.5

0.5

118

...

...

0.03

...

...

...

Udi

met

700

5515

175

...

...

43.

5..

...

...

.0.

060.

03..

...

.

Udi

met

710

5518

153

1.5

...

2.5

5..

...

...

.0.

070.

02..

...

.

Udi

met

720

5517

.914

.73

1.3

...

2.5

5..

...

...

.0.

030.

033

0.03

...

Uni

tem

p A

F2-1

DA

660

1210

2.7

6.5

...

42.

8..

...

...

.0.

040.

015

0.1

1.5

Ta

Was

palo

y58

19.5

13.5

4.3

...

...

1.3

3..

...

...

.0.

080.

006

...

...

Sour

ce:

Rep

rint

ed w

ith p

erm

issi

on fr

om S

tolo

ff, N

.S. 1

990.

ASM

Han

dboo

k Vo

l. 1.

Pro

pert

ies

and

Sele

ctio

n Ir

ons

Stee

ls a

nd H

igh

Perf

orm

ance

All

oys:

Wro

ught

and

P/M

Sup

eral

loys

. ASM

In

tern

atio

nal:

USA

, pp.

230

2–23

04.




TAB

LE 1

.5C

hem

ical

Com

posi

tion

of S

elec

ted

Cas

t N

i-B

ased

Sup

eral

loys

Allo

y

Com

posi

tion

(%

)

CC

rC

oM

oW

TaN

bA

lTi

Hf

Zr

BN

iO

ther

N-7

180.

0418

.5..

.3.

0..

...

.5

10.

50.

9..

...

...

.B

al18

.5 F

e

Ren

é 20

00.

0319

.012

.03.

2..

.3.

15.

10.

51.

0..

...

...

.B

al

IN-6

250.

0621

.5..

.8.

5..

...

.4.

00.

20.

2..

...

...

.B

al2.

5 Fe

IN-7

13C

0.12

12.5

...

4.2

...

...

2.0

6.1

0.8

...

0.10

0.01

2B

al..

.

IN-7

13L

C0.

0512

.0..

.4.

5..

...

.2.

05.

90.

6..

.0.

100.

01B

al..

.

IN-7

13 H

f (M

M 0

04)

0.05

12.0

...

4.5

...

...

2.0

5.9

0.6

1.3

0.10

0.01

Bal

...

IN-1

000.

1810

.015

.03.

0..

...

...

.5.

54.

7..

.0.

060.

014

Bal

1.0

V

IN-7

38C

0.17

16.0

8.5

1.75

2.6

1.75

0.9

3.4

3.4

...

0.10

0.01

Bal

...

IN-7

38L

C0.

1116

.08.

51.

752.

61.

750.

93.

43.

4..

.0.

040.

01B

al..

.

IN-7

920.

2112

.79.

02.

03.

93.

9..

.3.

24.

2..

.0.

100.

02B

al..

.

IN-9

390.

1522

.419

.0..

.2.

01.

41.

01.

93.

7..

.0.

100.

009

Bal

...

AM

1a..

.7

82

58

15.

01.

8..

...

...

.B

al

B-1

900

0.10

8.0

10.0

6.0

...

4.3

...

6.0

1.0

...

0.08

0.01

5B

al..

.

B-1

900

Hf

(MM

007

)0.

108.

010

.06.

0..

.4.

3..

.6.

01.

01.

50.

080.

015

Bal

...

B-1

910

0.10

10.0

10.0

3.0

...

7.0

...

6.0

1.0

...

0.10

0.01

5B

al..

.

MM

002

0.15

9.0

10.0

...

...

2.5

...

5.5

1.5

1.5

0.05

0.01

5B

al..

.

MA

R-M

200

0.15

9.0

10.0

...

12.5

...

1.8

5.0

2.0

...

0.05

0.01

5B

al..

.

MA

R-M

200

Hf

(MM

009

)0.

149.

010

.0..

.12

.5..

.1.

05.

02.

02.

0..

.0.

015

Bal

...

MA

R-M

246

0.15

9.0

10.0

2.5

10.0

1.5

...

5.5

1.5

...

0.05

0.01

5B

al..

.

MA

R-M

246

Hf

(MM

006

)0.

159.

010

.02.

510

.01.

5..

.5.

51.

51.

40.

050.

015

Bal

...

MA

R-M

247

(M

M 0

011)

0.16

8.5

10.0

0.65

10.0

3.0

...

5.6

1.0

1.4

0.04

0.01

5B

al..

.

CM

247

LC

0.07

8.1

9.3

0.5

9.5

3.0

...

5.6

0.7

1.4

0.01

0.01

5B

al..

.

PWA

148

0a..

.10

5..

.4

12..

.5.

01.

5..

...

...

.B

al

PWA

148

4a..

.5

102

69

...

5.6

...

0.1

...

...

Bal

3 R

e

Ren

é 41

0.08

19.0

10.5

9.5

...

...

...

1.7

3.2

...

0.01

0.00

5B

al..

.

Ren

é 77

0.08

15.0

18.5

5.2

...

...

...

4.25

3.5

...

...

0.01

5B

al..

.

Ren

é 80

0.17

14.0

9.5

4.0

4.0

...

...

3.0

5.0

...

0.03

0.01

5B

al..

. cont

inue

d



TAB

LE 1

.5

(con

tinu

ed)

Che

mic

al C

ompo

siti

on o

f Sel

ecte

d C

ast

Ni-

Bas

ed S

uper

allo

ys

Allo

y

Com

posi

tion

(%

)

CC

rC

oM

oW

TaN

bA

lTi

Hf

Zr

BN

iO

ther

Ren

é 80

Hf

0.15

14.0

9.5

4.0

4.0

...

...

3.0

4.7

0.8

0.01

0.01

5B

al..

.

Ren

é 10

00.

159.

515

.03.

0..

...

...

.5.

54.

2..

.0.

060.

015

Bal

1.0

V

Ren

é 12

5 H

f (M

M 0

05)

0.10

9.0

10.0

2.0

7.0

3.8

...

4.8

2.6

1.6

0.05

0.01

5B

al..

.

Ren

é N

4a..

.9

82

64

0.5

3.7

4.2

...

...

...

Bal

RR

200

0a..

.10

153

...

...

...

5.5

2.2

...

...

...

Bal

1 V

SRR

99a

...

85

...

103

...

5.5

2.2

...

...

...

Bal

Nim

ocas

t 75

0.12

20.0

...

...

...

...

...

...

0.5

...

...

...

Bal

...

Nim

ocas

t 80

0.05

19.5

...

...

...

...

...

1.4

2.3

...

...

...

Bal

1.5

Fe

Nim

ocas

t 90

0.06

19.5

18.0

...

...

...

...

1.4

2.4

...

...

...

Bal

1.5

Fe

Nim

ocas

t 95

0.07

19.5

18.0

...

...

...

...

2.0

2.9

...

0.02

0.01

5B

al..

.

Nim

ocas

t 100

0.20

11.0

20.0

5.0

...

...

...

5.0

1.5

...

0.03

0.01

5B

al..

.

Udi

met

500

0.08

18.5

16.5

3.5

...

...

...

3.0

3.0

...

...

0.00

6B

al..

.

Udi

met

700

0.08

14.3

14.5

4.3

...

...

...

4.25

3.5

...

0.02

0.01

5B

al..

.

Udi

met

710

0.13

18.0

15.0

3.0

1.5

...

...

2.5

5.0

...

0.08

...

Bal

...

C 1

300.

0421

.5..

.10

.0..

...

...

.0.

82.

6..

...

...

.B

al..

.

C 2

420.

3020

.010

.010

.3..

...

...

.0.

10.

2..

...

...

.B

al..

.

C 2

630.

0620

.020

.05.

9..

...

...

.0.

452.

15..

.0.

020.

001

Bal

...

C 1

023

0.15

15.5

10.0

8.0

...

...

...

4.2

3.6

...

...

0.00

6B

al..

.

Has

tello

y X

0.08

21.8

1.5

9.0

0.6

...

...

...

...

...

...

...

Bal

18.5

Fe,

0.5

M

n, 0

.3 S

i

Has

tello

y S

0.01

16.0

...

15.0

...

...

...

0.40

...

...

...

0.00

9B

al3.

0 Fe

, 0.0

2 L

a, 0

.65

Si,

0.55

Mn




Was

palo

y0.

0619

.012

.33.

8..

...

...

.1.

23.

0..

.0.

010.

005

Bal

0.45

Mn

NX

188

0.04

...

...

18.0

...

...

...

8.0

...

...

...

...

Bal

...

SEL

0.08

15.0

26.0

4.5

...

...

...

4.4

2.4

...

...

0.01

5B

al..

.

CM

SX-2

a..

.8.

04.

60.

68.

06.

0..

.5.

61.

0..

...

...

.B

al..

.

GM

R-2

350.

1515

.0..

.4.

8..

...

...

.3.

82.

0..

...

.0.

05B

al0.

3 M

n, 0

.4

Si, 1

1.0

Fe

CM

SX-3

a..

.8.

04.

60.

68.

06.

0..

.5.

61.

00.

10..

...

.B

al..

.

CM

SX-4

a..

.6.

49.

60.

66.

46.

5..

.5.

61.

00.

10..

...

.B

al3.

0 R

e

CM

SX-6

a..

.9.

95.

03.

0..

.2.

0..

.4.

84.

70.

05..

...

.B

al..

.

GM

R-2

350.

1515

.0..

.4.

8..

...

...

.3.

52.

5..

...

.0.

05B

al4.

5 Fe

SEL

-15

0.07

11.0

14.5

6.5

1.5

...

0.5

5.4

2.5

...

...

0.01

5B

al..

.

UD

M 5

60.

0216

.05.

01.

56.

0..

...

.4.

52.

0..

.0.

030.

070

Bal

0.5

V

M-2

20.

135.

7..

.2.

011

.03.

0..

.6.

3..

...

.0.

60..

.B

al..

.

IN-7

310.

189.

510

.02.

5..

...

...

.5.

54.

6..

.0.

060.

015

Bal

1.0

V

MA

R-M

421

0.14

15.8

9.5

2.0

3.8

...

...

4.3

1.8

...

0.05

0.01

5B

al..

.

MA

R-M

432

0.15

15.5

20.0

...

3.0

2.0

2.0

2.8

4.3

...

0.05

0.01

5B

al..

.

MC

-102

0.04

20.0

...

6.0

2.5

0.6

6.0

...

...

...

...

...

Bal

0.25

Si,

0.30

M

n

Nim

ocas

t 242

0.34

20.5

10.0

10.5

...

...

...

0.2

0.3

...

...

...

Bal

1.0

Fe, 0

.3

Mn,

0.3

Si

Nim

ocas

t 263

0.06

20.0

20.0

5.8

...

...

...

0.5

2.2

...

0.04

0.00

8B

al0.

5 Fe

, 0.5

M

n

Sour

ce:

Rep

rint

ed w

ith p

erm

issi

on f

rom

Har

ris,

K.,

Eri

ckso

n, G

.L.,

Schw

er,

R.E

. 19

90.

ASM

Han

dboo

k Vo

l. 1.

Pro

pert

ies

and

Sele

ctio

n Ir

ons

Stee

ls a

nd H

igh

Perf

orm

ance

All

oys:

Po

lycr

ysta

llin

e C

ast S

uper

allo

ys &

Dir

ecti

onal

ly S

olid

i�ed

and

Sin

gle

Cry

stal

All

oys.

ASM

Int

erna

tiona

l: U

SA. p

p. 2

386–

2389

& 2

427–

2429

.a

Sing

le c

ryst

al.



TAB

LE 1

.6C

hem

ical

Com

posi

tion

of S

elec

ted

Cas

t C

o-B

ased

Sup

eral

loys

Allo

y

Com

posi

tion

(%

)

CC

rN

iW

TaN

bM

oTi

BZ

rFe

Co

Oth

er

HS-

21 (

MO

D V

italli

um)

0.25

27.0

3.0

...

...

...

5.0

...

...

...

1.0

Bal

HS-

31 (

X-4

0)0.

5025

.010

.07.

5..

...

...

...

...

.0.

171.

5B

al0.

4 Si

HS-

25 (

L-6

05)

0.10

20.0

10.0

15.0

...

...

...

...

...

...

...

Bal

ML

-170

00.

225

.0..

.15

.0..

...

...

...

.0.

4..

...

.B

al

WI-

520.

4221

.01.

0 m

ax11

.0..

.2.

0..

...

...

...

.2.

0B

al

MA

R-M

302

0.85

21.5

...

10.0

9.0

...

...

0.2

0.00

5..

.1.

5 m

axB

al

MA

R-M

322

1.0

21.5

...

9.0

4.5

...

...

0.75

...

2.25

0.75

Bal

MA

R-M

509

0.60

24.0

10.0

7.0

7.5

...

...

0.2

...

...

1.0

Bal

AiR

esis

t 13

0.45

21.0

...

11.0

...

2.0

...

...

...

...

2.5

max

Bal

3.4

A1,

0.1

Y

AiR

esis

t 215

0.35

19.0

0.5

4.5

7.5

...

...

...

...

0.13

...

Bal

4.3

A1,

0.1

Y

F 75

0.25

28.0

1.0

max

...

...

...

5.5

...

...

...

...

Bal

FSX

-414

0.25

29.5

10.5

7.0

...

...

...

...

0.01

2..

.2.

0 m

axB

al

X-4

50.

2525

.510

.57.

0..

...

...

...

.0.

010

...

2.0

max

Bal

Sour

ce:

Rep

rint

ed w

ith p

erm

issi

on f

rom

Har

ris,

K.,

Eri

ckso

n, G

.L.,

Schw

er,

R.E

. 19

90.

ASM

Han

dboo

k Vo

l. 1.

Pro

pert

ies

and

Sele

ctio

n Ir

ons

Stee

ls a

nd H

igh

Perf

orm

ance

All

oys:

Po

lycr

ysta

llin

e C

ast S

uper

allo

ys. A

SM I

nter

natio

nal:

USA

, p. 2

390.




TAB

LE 1

.7M

echa

nica

l Pro

pert

ies

of S

elec

ted

Supe

rallo

ys

Allo

yFo

rm

Ult

imat

e Te

nsile

Str

engt

h at

21°C

(70

°F)

540°

C

(100

0°F)

650°

C

(120

0°F)

760°

C

(140

0°F)

870°

C

(160

0°F)

Mpa

ksi

Mpa

ksi

Mpa

ksi

Mpa

ksi

Mpa

ksi

Con

diti

on o

f Tes

t M

ater

iala

Nic

kel-

Bas

edA

stro

loy

Bar

1415

205

1240

180

1310

190

1160

168

775

112

1095

°C (

2000

°F)/

4 h/

OQ

+ 8

70°C

(16

00°F

)/8

h/A

C +

980

°C (

1800

°F)/

4 h/

AC

+ 6

50°C

(12

00°F

) 24

h/

AC

+ 7

60°C

(14

00°F

)/8

h/A

C

Cab

ot 2

14..

.91

513

371

510

467

598

560

8444

064

1120

°C (

2050

°F)

D-9

79B

ar14

1020

412

9518

811

0516

072

010

434

550

1040

°C (

1900

°F)/

1 h/

OQ

+ 8

45°C

(15

50°F

)/6

h/A

C +

705

°C(1

300°

F)/1

6 h/

AC

Has

tello

y C

-22

Shee

t80

011

662

591

585

8552

576

...

...

1120

°C (

2050

°F)/

RQ

Has

tello

y G

-30

Shee

t69

010

049

071

...

...

...

...

...

...

1175

°C (

2150

°F)/

RA

C-W

Q

Has

tello

y S

Bar

845

130

775

112

720

105

575

8434

050

1065

°C (

1950

°F)/

AC

Has

tello

y X

Shee

t78

511

465

094

570

8343

563

255

3711

75°C

(21

50°F

)/1

h/R

AC

Hay

nes

230

...

870

126

720

105

675

9857

584

385

5612

30°C

(22

50°F

)/A

C

Inco

nel 5

87b

Bar

1180

171

1035

150

1005

146

830

120

525

76..

.

Inco

nel 5

97b

Bar

1220

177

1140

165

1060

154

930

135

...

...

...

Inco

nel 6

00B

ar66

096

560

8145

065

260

3814

020

1120

°C (

2050

°F)/

2 h/

AC

Inco

nel 6

01Sh

eet

740

107

725

105

525

7629

042

160

2311

50°C

(21

00°F

)/2

h/A

C

ncon

el 6

17B

ar74

010

758

084

565

8244

064

275

4011

75°C

(21

50°F

)/A

C

Inco

nel 6

17Sh

eet

770

112

590

8659

086

470

6831

045

1175

°C (

2150

°F)/

0.2

h/A

C

Inco

nel 6

25B

ar96

514

091

013

283

512

155

080

275

4011

50°C

(21

00°F

)/1

h/W

Q

Inco

nel 7

06B

ar13

1019

011

4516

610

3515

072

510

5..

...

.98

0°C

(18

00°F

)/1

h/A

C +

845

°C (

1550

°F)/

3 h/

AC

+ 7

20°C

(13

25°F

)/8

h/FC

+62

0°C

(115

0°F)

/8 h

/AC

Inco

nel 7

18B

ar14

3520

812

7518

512

2817

895

013

834

049

980°

C (1

800°

F)/1

h/A

C +

720

°C (1

325°

F)/8

h/

FC +

620

°C (

1150

°F)/

18 h

/AC

cont

inue

d



TAB

LE 1

.7

(con

tinu

ed)

Mec

hani

cal P

rope

rtie

s of

Sel

ecte

d Su

pera

lloys

Allo

yFo

rm

Ult

imat

e Te

nsile

Str

engt

h at

21°C

(70

°F)

540°

C

(100

0°F)

650°

C

(120

0°F)

760°

C

(140

0°F)

870°

C

(160

0°F)

Mpa

ksi

Mpa

ksi

Mpa

ksi

Mpa

ksi

Mpa

ksi

Con

diti

on o

f Tes

t M

ater

iala

Inco

nel 7

18

Dir

ect A

geB

ar15

3022

213

5019

612

3517

9..

...

...

...

.73

5°C

(13

25°F

)/8

h/SC

+ 6

20°C

(11

50°F

)/8

h/A

C

Inco

nel 7

18

Supe

rB

ar13

5019

612

0017

411

3016

4..

...

...

...

.92

5°C

(17

00°F

)/1

h/A

C +

735

°C (

1325

°F)/

8 h/

SC +

620

°C (

1150

°F)/

8 h/

AC

Inco

nel X

750

Bar

1200

174

1050

152

940

136

...

...

...

...

1150

°C (

2100

°F)/

2 h/

AC

+ 8

45°C

(15

50°F

)/24

h/

AC

+ 7

05°C

(13

00°F

)/20

h/A

C

M-2

52B

ar12

4018

012

3017

811

6016

894

513

751

074

1040

°C (

1900

°F)/

4 h/

AC

+ 7

60°C

(14

00°F

)/16

h/A

C

Nim

onic

75

Bar

745

108

675

9854

078

310

4515

022

1050

°F (

1925

°F)/

1 h/

AC

108

0°C

(19

75°F

)/8

h/A

C +

705

°C (

1300

°F)/

16 h

/AC

Nim

onic

90

Bar

1235

179

1075

156

940

136

655

9533

048

1080

°C (

1975

°F)/

8 h/

AC

+ 7

05°C

(13

00°F

)/16

h/A

C

Nim

onic

105

Bar

1180

171

1130

164

1095

159

930

135

660

9611

50°C

(21

00°F

)/4

h/A

C +

106

0°C

(19

40°F

)/16

h/

AC

+ 8

50°C

(15

60°F

)/16

h/A

C

Nim

onic

115

Bar

1240

180

1090

158

1125

163

1085

157

830

120

1190

°C (

2175

°F)/

1.5

h/A

C +

110

0°C

(20

10°F

)/6

h/A

C

Nim

onic

263

Shee

t97

014

180

011

677

011

265

094

280

4011

50°C

(21

00°F

)/0.

2 h/

WQ

+ 8

00°C

(14

70°F

)/8

h/A

C

Nim

onic

942

bB

ar14

0520

413

0018

912

4018

090

013

1..

...

...

.

Nim

onic

PE

.11b

Bar

1080

157

1000

145

940

136

760

110

...

...

...

Nim

onic

PE

.16

Bar

885

128

740

107

660

9651

074

215

3110

40°C

(190

0°F

)/4

h/A

C +

800

°C (1

470°

F)/

2 h/

AC

+ 7

00°C

(129

0°F

)/16

h/A

CN

imon

ic P

K.3

3Sh

eet

1180

171

1000

145

1000

145

885

128

510

7411

00–1

115°

C (

2010

–204

0°F)

/0.2

5 h/

AC

+ 8

50°C

(1

500°

F)/4

h/A

C

Pyro

met

860

bB

ar12

9518

812

5518

211

1016

191

013

2..

...

.10

95°C

(20

00°F

)/2

h/W

Q +

830

°C (

1525

°F)/

2 h/

AC

+ 7

60°C

(14

00°F

)/24

h/A

C

Ren

é 41

Bar

1420

206

1400

203

1340

194

1105

160

620

9010

65°C

(19

50°F

)/4

h/A

C +

760

°C (

1400

°F)/

16 h

/AC

Ren

é 95

Bar

1620

235

1550

224

1460

212

1170

170

...

...

900°

C (

1650

°F)/

24 h

+ 1

105°

C (

2025

°F)/

1 h/

OQ

+ 7

30°C

(13

50°F

)/64

h/A

CFor cataloging purposes only



Udi

met

400

bB

ar13

1019

011

8517

2..

...

...

...

...

...

...

.

Udi

met

500

Bar

1310

190

1240

180

1215

176

1040

151

640

9310

80°C

(19

75°F

)/4

h/A

C +

845

°C (

1550

°F)/

24 h

/A

C +

760

°C (

1400

°F)/

16 h

/AC

Udi

met

520

Bar

1310

190

1240

180

1175

170

725

105

515

7511

05°C

(20

25°F

)/4

h/A

C +

845

°C (

1550

°F)/

24 h

/A

C +

760

°C (

1400

°F)/

16 h

/AC

Udi

met

630

bB

ar15

2022

013

8020

012

7518

596

514

0..

...

...

.

Udi

met

700

Bar

1410

204

1275

185

1240

180

1035

150

690

100

1175

°C (

2150

°F)/

4 h/

AC

+ 1

080°

C (

1975

°F)/

4 h/

AC

+ 8

45°C

(15

50°F

)/24

h/A

C +

760

°C (

1400

°F)/

16 h

/A

C

Udi

met

710

Bar

1185

172

1150

167

1290

187

1020

148

705

102

1175

°C (

2150

°F)/

4 h/

AC

+ 1

080°

C (

1975

°F)/

4 h/

AC

+ 8

45°C

(15

50°F

)/24

h/A

C +

760

°C

(140

0°F)

/16

h/A

C

Udi

met

720

Bar

1570

228

...

...

1455

211

1455

211

1150

167

1115

°C (

2035

°F)/

2 h/

AC

+ 1

080°

C (

1975

°F)/

4 h/

OQ

+ 6

50°C

(12

00°F

)/24

h/A

C +

760

°C (

1400

°F)/

8 h/

AC

Uni

tem

p A

F2-1

DA

6B

ar15

6022

614

8021

514

0020

312

9018

7..

...

.11

50°C

(21

00°F

)/4

h/A

C +

760

°C (

1400

°F)/

16 h

/AC

Was

palo

yB

ar12

7518

511

7017

011

1516

265

094

275

4010

80°C

(19

75°F

)/4

h/A

C +

845

°C (

1550

°F)/

24 h

/AC

+

Iron

-Bas

edA

-286

Bar

1005

146

905

131

720

104

440

64..

...

.98

0°C

(18

00°F

)/1

h/O

Q +

720

°C (

1325

°F)/

16 h

/AC

Allo

y 90

1B

ar12

0517

510

3014

996

013

972

510

5..

...

.10

95°C

(20

00°F

)/2

h/W

Q +

790

°C (

1450

°F)/

2 h/

AC

+ 7

20°C

(13

25°F

)/24

h/A

C

Dis

calo

yB

ar10

0014

586

512

572

010

448

570

...

...

1010

°C (

1850

°F)/

2 h/

OQ

+ 7

30°C

(13

50°F

)/20

h/

AC

+ 6

50°C

(12

00°F

)/20

h/A

C

Hay

nes

556

Shee

t81

511

864

593

590

8547

069

330

4811

75°C

(21

50°F

)/A

C

Inco

loy

800b

Bar

595

8651

074

405

5923

534

...

...

...

Inco

loy

801b

Bar

785

114

660

9654

078

325

47..

...

...

.

Inco

loy

802b

Bar

690

100

600

8752

576

400

5819

528

...

Inco

loy

807b

Bar

655

9547

068

440

6435

051

220

32..

.

Inco

loy

825cd

...

690

100

~590

~86

~470

~68

~275

~40

~140

~20

...

Inco

loy

903

Bar

1310

190

...

...

1000

145

...

...

...

...

845°

C (

1550

°F)/

1 h/

WQ

+ 7

20°C

(13

25°F

)/8

h/FC

+6

20°C

(11

50°F

)/8

h/A

C

cont

inue

d



TAB

LE 1

.7

(con

tinu

ed)

Mec

hani

cal P

rope

rtie

s of

Sel

ecte

d Su

pera

lloys

Allo

yFo

rm

Ult

imat

e Te

nsile

Str

engt

h at

21°C

(70

°F)

540°

C

(100

0°F)

650°

C

(120

0°F)

760°

C

(140

0°F)

870°

C

(160

0°F)

Mpa

ksi

Mpa

ksi

Mpa

ksi

Mpa

ksi

Mpa

ksi

Con

diti

on o

f Tes

t M

ater

iala

Inco

loy

907ce

...

~136

5~1

98~1

205

~175

~103

5~1

50~6

55~9

5..

...

...

.

Inco

loy

909

Bar

1310

190

1160

168

1025

149

615

89..

...

.98

0°C

(18

00°F

)/1

h/W

Q +

720

°C (

1325

°F)/

8 h/

FC +

620

°C (

1150

°F)/

8 h/

AC

N-1

55B

ar81

511

865

094

545

7942

862

260

3811

75°C

(21

50°F

)/1

h/W

Q +

815

°C (

1500

°F)/

4 h/

AC

V-5

7B

ar11

7017

010

0014

589

513

062

090

...

...

980°

C (

1800

°F)/

2–4

h/O

Q +

730

°C (

1350

°F)/

16 h

/AC

19-9

DL

f..

.81

511

861

589

517

75..

...

...

...

...

.

16-2

5-6f

...

980

142

...

...

620

9041

560

...

...

...

Cob

alt-

Bas

ed

Air

Res

ist 2

13g

...

1120

162

...

...

960

139

485

7031

546

...

Elg

iloyg

...

690d –

2480

h10

0d –36

0h..

...

...

...

...

...

...

...

...

.

Hay

nes

188

Shee

t96

013

974

010

771

010

363

592

420

6111

75°C

(21

50°F

)/1

h/R

AC

L-6

05Sh

eet

1005

146

800

116

710

103

455

6632

547

1230

°C (

2250

°F)/

1 h/

RA

C

MA

R-M

918

Shee

t89

513

0..

...

...

...

...

...

...

...

.11

90°C

(21

75°F

)/4

h/A

C




MP3

5NB

ar20

2529

4..

...

...

...

...

...

...

...

.53

% C

W +

565

°C (

1050

°F)/

4 h/

AC

MP1

59B

ar18

9527

515

6522

715

4022

3..

...

...

...

.48

% C

W +

665

°C (

1225

°F)/

4 h/

AC

Stel

lite

6Bg

Shee

t10

1014

6..

...

...

...

...

...

.38

556

2 m

m (

0.06

3 in

.) s

heet

hea

t tre

ated

at 1

232°

C (

2250

°F)

and

RA

C

Hay

nes

150g

...

925

134

...

...

325i

47..

...

.15

5j22

.7..

.

Sour

ce:

Rep

rint

ed w

ith p

erm

issi

on f

rom

Sto

loff

, N.S

. 199

0. A

SM H

andb

ook

Vol.

1. P

rope

rtie

s an

d Se

lect

ion

Iron

s St

eels

and

Hig

h Pe

rfor

man

ce A

lloy

s: W

roug

ht a

nd P

/M S

uper

allo

ys.

ASM

Int

erna

tiona

l: U

SA, p

p. 2

317–

2321

.a

OQ

, oil

quen

ch; A

C, a

ir c

ool;

RQ

, rap

id q

uenc

h; R

AC

-WQ

, rap

id a

ir c

ool-

wat

er q

uenc

h; F

C, f

urna

ce c

ool;

SC, s

low

coo

l; C

W, c

old

wor

ked.

b D

ata

refe

rred

fro

m H

igh-

Tem

pera

ture

Hig

h-St

reng

th N

icke

l Bas

e A

lloy

s, I

nco

Allo

ys I

nter

natio

nal L

td.,

dist

ribu

ted

by N

icke

l Dev

elop

men

t Ins

titut

ec

Dat

a re

ffer

ed f

rom

Pro

duct

Han

dboo

k, P

ublic

atio

n 1A

1-38

, Inc

o A

lloys

Int

erna

tiona

l, In

c., 1

988

d A

nnea

led.

e Pr

ecip

itatio

n ha

rden

ed.

f D

ata

refe

rred

fro

m M

ater

ials

Sel

ecto

r, 19

88, P

ento

n, 1

987.

g D

ata

refe

rred

fro

m F

.R. M

orra

l, E

d., W

roug

ht s

uper

allo

ys, i

n P

rope

rtie

s an

d Se

lect

ion:

Sta

inle

ss S

teel

s, T

ool M

ater

ials

and

Spe

cial

-Pur

pose

Met

als,

Vol

ume

3, 9

th e

d., M

etal

Han

dboo

k,

Am

eric

an S

ocie

ty f

or M

etal

s, 1

980.

h W

ork

stre

ngth

ened

and

age

d.i

At 7

00°C

(12

90°F

).j

At 9

00°C

(16

50°F

).



1.4.1 CASTING–MELTING PRACTICE

All cast, wrought, and powder metallurgy products need initial melting to make their raw material in an ingot or in an electrode form. Although several melting methods exist today, their purpose is focused on two primary areas: metal cleanliness and structural uniformity.

1.4.1.1 Melt ProcessingThe schematic view of the main melting of superalloys is given in Figure 1.15. Double or sometimes even triple melting is applied to achieve cleanliness and re�ned uniformity. The main unwanted ele-ments are Bi, Pb, S, and so on that usually form as oxides, nitrides, and sul�des. Re�ned uniformity covers both chemical composition and grain solidi�cation structure.

Although dif�cult and expensive, vacuum, instead of air, is very bene�cial in the melting prac-tice. Oxygen, hydrogen, and nitrogen in the charge are greatly reduced. Some detrimental and vola-tile elements such as Bi, Pb, and Se found in raw materials are easily degassed. Furthermore, easily oxidized reactive elements such as Al, Ti, and Hf can be effectively added into the composition in much higher amount in vacuum. The main unreactive elements in superalloys are Co, Ni, Fe, Cr, Mo, W, and Ta (Tien and Caul�eld 1989).

As can be seen in Figure 1.15, VIM is used in converting the raw materials and scrap metal to ingots or �rst producing electrodes for successive remelting. Scrap metal may come from various sources such as rejected parts, chips from machining operations, and so on. Generally, used parts are not acceptable for direct VIM melting due to their contamination such as sulfur from service history.

Although some of the cobalt- and iron–nickel-based superalloys, which do not contain reactive elements, can be melted in air instead of vacuum, these alloys are usually remelted with various techniques like vacuum arc remelting (VAR), electroslag remelting (ESR), and so on. The simpli�ed VIM process schematic is given in Figure 1.16.

As seen in Figure 1.16, alloying elements are melted inductively by hollow copper coils through which water and electric current pass, then poured to the mold via a tundish to solidify as ingot or

Scrap and recycledalloys

(machining chips,misproduced

parts, used parts,etc.)

Primary melting

Raw materials(Fe, Ni, Co, Cr,

etc.)

Resizing

ResizingIngot

Electrode

Vacuum inductionmelting(VIM,

atomizationprocesses for

powdermetallurgy)

Vacuum inductionmelting(VIM)

Investmentcasting

Vacuum arcremelting

(VAR)

Electroslagremelting

(ESR)

Secondary melting Triple melting

Vacuum arcdouble electrode

remelting(VADER)

Vacuum arcremelting

(VAR)

FIGURE 1.15 Schematic view of melting practices applied to the superalloys.




electrode under vacuum. The circulating water acts as an electric carrier and cooling agent. During pouring, usually ceramic �lters are used to eliminate relatively large oxide and nitride inclusions (Campbell 2006). As is natural in the induction melting process, the alloy subjected to melting is continuously stirred by magnetic �elds; hence, the alloy chemistry is homogenized. The main disadvantage of this stirring movement is that oxide particles with their low densities cannot come out or �oat on the surface (Tien and Caul�eld 1989). Not all charging elements are placed in the crucible at the beginning of the process; �rst, unreactive elements such as Co and Ni are melted and degassed, then reactive elements such as Zr, Hf, and Al are added via an air lock which can be seen in Figure 1.16 (Tien and Caul�eld 1989). Reactive elements in the charge are usually obtained from scrap metal rather than raw material in commercial melting procedure.

The ingots obtained are usually heavily segregated, containing shrinkage, coarse, and nonuni-form grains. This structure is not important for subsequent investment casting but it is not suitable for wrought process, and hence, a second melting is required. The most widely used methods for remelting are VAR and ESR (Campbell 2006).

In VAR, an arc is created between ingots, which acts as an electrode, and water-cooled crucible bottom, as melting continues, the melting electrode is shortened and molten metal in the crucible rises and solidi�es. The shortened electrode is pulled up with a mechanism; thus the gap between the molten pool and the electrode is kept constant (Tien and Caul�eld 1989). The schematic view of VAR process is given in Figure 1.17.

The VIM electrode turns into an ingot with greatly improved chemical and physical homogene-ity by the VAR process. Grain size and solidi�cation microstructure are re�ned in the process and the remaining volatile components like bismuth and lead are further reduced. As can be seen in Figure 1.17, the VAR process creates semi-spherical U-shaped liquid pool. The size and geometry of this pool is closely related with the process parameters and affects the �nal quality of the ingot.

Charge vacuumlock cover

Vacuum chamber

Launder

Tundish

Inductioncoil

Mold

Solidifiedmetal

Ceramicfilter

Crucible

Molten metal

Water + electric line

Water cooler

Power supply

Vacuum line

Vacuumsystem

FIGURE 1.16 (See color insert.) Schematic view of vacuum induction melting process.



The �oating oxides and other impurities in the molten pool are swept to the outer diameter of the ingot by an arc, called shelf. After vacuum application, often helium gas with small partial pressure is introduced into the vacuum chamber to improve the ef�ciency of the arc.

VAR is susceptible to electrode fall and shelf material collapse which are sources of some defects during melting process. The main reasons for electrode fall are

• VIM electrode is of too much shrinkage cavities, internal porosity and cracks• Electrode cracking due to severe cooling conditions during VAR• Preferential melting due to melting point difference between interdendritic material and

dendrite spine (Tien and Caul�eld 1989)

In the event of top portion of the shelf material detaching from the crucible wall during melting, the solidi�ed metal may also collapse and fall into the molten pool.

Both electrode fall and shelf material collapse can cause inclusions and inhomogeneities in the �nal ingot. These defects can be reduced by using sound electrodes and careful application of pro-cess parameters. Sometimes, VIM + VAR electrodes are used instead of VIM electrodes to increase electrode soundness. Advanced monitoring and closed loop control systems are used frequently to adjust key parameters and maintain best melting throughout the process.

In ESR method, the slag is used as protective environment instead of vacuum. The electrode is placed in water cooled crucible similar to VAR and covered with a protective slag. Although vari-ous materials can be used as slag, the most widely used slag materials are CaF2, CaO, Al2O3, MgF2 (Tien and Caul�eld 1989). A schematic view of ESR process is given in Figure 1.18.

Ram (drawback mechanism)

ElectrodeArc gap (betweenelectrode and molten pool)

Molten pool

Power supplyVacuum pumpWater coolingCopper crucibleWater jacket

Solidified metal

Dendrites

FIGURE 1.17 (See color insert.) Schematic view of vacuum arc remelting process.




The required heat is supplied from electrical resistance of slag, the molten droplets from elec-trode pass through molten slag and solidify on water cooled crucible bottom. During the passing of droplets, much of the unwanted contaminants are caught by slag. As the molten metal pool surface is completely covered with slag, the heat transfer from the surface decreases thus shelf thickness is thinner than it would be at VAR. Also, the turbulence within the molten metal is far lower than that during VAR due to lack of arc, therefore, inclusions not caught by slag during passing would easily �oat on the surface and get caught by slag.

As in the VAR method, falling electrode is also possible in ESR. The fallen metal may drag a large amount of slag to the molten pool and this residue may solidify in the ingot as a defect source. Volatile tramp elements such as Bi and Pb are not subjected to further re�ning by degassing as in VAR, as they are not in a vacuum environment. Therefore, the maximum permissible limit of these elements should be met by preceding melting practice. The molten metal pool shape is different from that during VAR and deeply grooved V in ESR, because of reduced heat transfer by slag. As a result of this solidi�cation front shape, the center of the ESR ingots have an unwanted coarse dendritic structure. To compensate for this situation, smaller diameter ingots are produced in ESR process. As in VAR, advanced monitoring and closed loop control systems are frequently used to improve the process.

The ESR process produces cleaner alloys without an expensive and complex vacuum environ-ment. However, the VAR process is capable of producing larger ingots with more reduced segrega-tion defects. As a result, the requirement of large ingots used in massive parts produced by forging is met by triple (VIM + ESR + VAR) melting (Tien and Caul�eld 1989).

Ram (drawback mechanism)

Electrode

Molten metal droplets

Molten slag

Molten pool

Power supplyWater coolingCopper crucibleWater jacket

Solidified metal

Dendrites

FIGURE 1.18 (See color insert.) Schematic view of electrical slag remelting process.



1.4.1.2 Investment CastingAlthough, the primitive form of investment casting goes back to the Bronze Age, the �rst industrial usage is dentistry at the beginning of the 1900s. The �rst application of modern investment casting in recent times can be seen in the production of cobalt-based turbine blades in the 1940s. Nearly 35 million cobalt-based vitallium turbine blades were produced by investment casting until the end of World War II. Turbine blades were cast using electric arc furnaces at atmospheric pressure in those years. Vacuum melting was performed then vacuum casting was applied to the superalloys in the 1950s (Sims et al. 1987; Tien and Caul�eld 1989). Vacuum technology made the production of high-strength nickel-based superalloys possible, thanks to the ability of inclusion of higher amount of reactive elements. Besides, vacuum casting is also very effective in reducing defects such as gas cavities and un�lling metal in complex geometries.

Another important development in the investment-casting area is the use of a self-supporting shell instead of a �ask that surrounds the mold since the late 1950s. The alumina zirconia-based shell materials instead of silica magnesium oxide mixture bring various advantages. The thin-ner shell structure allows faster cooling of the cast and as a result re�ned microstructure. The shell clusters, formed through gathering a number of self-supporting shells, greatly increase productivity. The resistance to thermal gradients of shell makes directional solidi�cation and single crystal alloy production possible.

The most important developments in the 1960s and the 1980s are DS and SX, respectively. The �rst known industrial application of SX turbine blade in aerospace industry was applied to the Pratt & Whitney JT9D-7R4 civil aircraft gas turbine in 1982 (Gell 1985). Another very important advancement in investment casting is the production of ceramic cores, which give possibility to obtain complex internal cooling geometries. When this geometry is applied to the turbine blades and vanes, higher turbine entry temperature (TIT) is possible.

A schematic view of the main production steps of investment-casting process is given in Figure 1.19. Although turbine blades and vanes are frequently produced by DS and SX tech-niques, many other parts of gas turbine engines such as frames, diffuser case are still produced by PC today. Therefore, all of the investment-casting techniques, including polycrystalline casting, are still widely used.

The �rst step is model making. As in other conventional casting techniques, determining model dimension should take into account various factors such as shrinkage values of alloys, mold, and model. Natural and synthetic waxes, resins, and �llers are generally used as pattern material. A good pattern material should have speci�c properties such as low expansion, low contraction, enough mechanical strength, low ash content, dimensional stability, and so on (Tien and Caul�eld 1989).

Besides patterns, cores should also be used when producing parts that contain internal cavities like turbine blade cooling passages and holes. Cores are placed in the die before wax injection thus injected wax �lls the remaining volume of the die. During placement of cores and injection of wax, care must be taken to prevent core breakage or core shift. Less frequently, quartz tubes can be used instead of cores at production of cooling holes. Like patterns, the core itself should also have enough strength and its expansion and contraction characteristics should be in conformity with the pattern, mold, and die. Unlike wax, the core stays in the shell until the end of the casting process and thus has to come into contact with hot molten metal for a considerable amount of time. Therefore, the core material should also be unreactive to the molten metal and should be easily removed after casting without any damage to the alloy. Generally, alumina- and silica-based materials are used as core. Turbine blade and core for producing turbine blades’ complex cooling schemes can be seen in Figure 1.20.

Although a preheating process is applied to the mold, temperature differences still exist between molten metal and core in PC. Therefore, the cores are also subject to thermal shocks and must not crack during pouring of hot molten metal. This temperature difference is minimized in DS and SX process but the core stays much longer �rst in molten then hot solidi�ed metal when compared with




Die making by usingmold

Wax clustering

Casting Cutting and finishing Investment-castedturbine blade

Cluster dying withslurry

Stuccoing by powderinjection Dewaxing

Die waxing Wax mold

FIGURE 1.19 Main production steps of investment-casting process.



PC. Hence, the chemical stability of the core is very critical to avoid deleterious surface reactions between the core and the alloy.

After placing the core (if it exists) in the die, the wax is injected to produce a pattern. Several patterns are brought together with their gates, runner, and pouring cap to form clusters. Clustering is a very ef�cient way to enhance producibility, reliability, and saving a considerable amount of time. The schematic view of a typical cluster is in Figure 1.21.

FIGURE 1.20 Turbine blade (outer wall removed; left) and core (right) with complex shape. (Reprinted with permission of Donachie, M.J., Donachie, S.J. 2002. Superalloys A Technical Guide, p. 82, ASM International: USA.)

FIGURE 1.21 Schematic view of a cluster.




Clusters are subjected to successive dipping in the slurry and are coated with dry powder called stucco to obtain a shell. The number of this repeated process typically is 5–15 to obtain enough thickness and strength. Shell thickness is an important parameter to control the cooling rate. The �rst dipping of a cluster and the following stucco have special importance as the molten metal directly contacts this surface. This layer is called a “facecoat” and governs the �nal product sur-face roughness and shape that may include shallow or small features. Some nucleating agents like cobalt aluminate are also added to facecoat for grain re�nement. The remaining layers only act as support to prevent cracking of the mold during casting and govern heat transfer rate. Computer controlled and automated facilities seen in Figure 1.22 are frequently used to obtain reliable, repeatable shells.

Typical shell materials and their maximum service temperatures are tabulated in Table 1.8.The cluster is dried and dewaxed. Both processes take a long time to ensure the mold’s dimen-

sional and mechanical stability. Naturally, dewaxing is performed under a higher temperature than drying. The cluster is heated before casting to control solidi�cation rate and is usually placed in an insulator to reduce heat transfer before casting.

FIGURE 1.22 Automated dipping process. (Reprinted with permission of Donachie, M.J., Donachie, S.J. 2002. Superalloys A Technical Guide, p. 84, ASM International: USA.)

TABLE 1.8Typical Shell Materials and Their Maximum Application Temperatures

MaterialMaximum Application

Temperature (Approximate)

Al2O3 1840°C

Al2O3–SiO2 (aluminasilicate with varying ratios, typically 20–50 SiO2%)

1560°C

SiO2 1680°C

ZrSiO4 1650°C

Source: Adapted from Sims, C.T. Stoloff, N.S. Hagel, W.C. 1987. Superalloys II. John Wiley & Sons: USA; Stefenescu, D.M. 1998. ASM Handbook Vol. 15. Casting: Investment casting. ASM International: USA.)



The remelting is very similar to the primary VIM melting. The required amount of material is cut from the ingot, placed in a crucible, inductively melted and poured to the preheated mold then cooled to ambient temperature. With a few exceptions of some cobalt-based alloys, both melting and pouring are done in vacuum to obtain the best results. Unlike primary melting, now the unwanted chemical reactions between the crucible and the molten metal are very limited as both melting and pouring time are short (Tien and Caul�eld 1989).

The molten metal temperature is approximately 80–160°C higher than the liquidus point (super heat) and this difference is a critical parameter that affects the solidi�cation rate, grain size, grain orientation, and location of microshrinkage and hence is measured with optical pyrometers continu-ously throughout the process. The preheating temperature for the mold depends on various casting parameters such as shape of the part, desired grain sizes, and so on. In conventional polycrystalline investment casting, the mold temperature is kept well below the liquidus of the metal (usually at ~850–1250°C); in directional or single crystal investment casting, however, the mold temperature needs to be above the liquidus (Sims et al. 1987).

After a suitable temperature range for both mold and molten metal is established, the molten metal is rapidly poured from the crucible to the mold. This process is also critical: pouring rate, exact positioning of the mold and crucible all affect casting quality. Therefore, to ensure high qual-ity cast, all processes including melting and casting are widely done via automated control systems as in mold production.

As mentioned previously, grain size is an important parameter of strength. Coarse grains give rise to better creep and rupture properties especially at high temperatures while a �ne-grained microstructure brings better fatigue and tensile properties especially at equicohesive temperatures. The casting volume signi�cantly affects the �nal grain size; low volume parts such as turbine blades and vanes can be cooled rapidly thus the desired grain sizes are easily obtained. But turbine disks are of a larger volume that need longer time to complete the solidi�cation especially in large sec-tions like hub. As a result, unwanted coarse grains form.

As the service temperature for turbine disks is signi�cantly lower than that of the turbine blade and vanes, the disk needs a small grain microstructure to achieve better low cycle fatigue and tensile proper-ties. Special alloys like Grainex® and Microcast® are designed for turbine wheels to achieve �ne micro-structures, with typically ASTM 2-5 grain sizes. In processing of Grainex mold agitation is used; for Microcast, low degree of superheat is applied with turbulent �ow during casting to promote nucleation.

Although melting and many other aspects of investment casting are nearly the same in conven-tional PC, DS, or SX techniques, mold design and cooling scheme are considerably different. The schematic views of PC, DS, and SX mold design are given in Figure 1.23.

Vacuum environment

Induction coil

Refracter wall

Molten metal

Solidified metalHelical mold selector

Water cooled copperchill block

Drawback mechanism

FIGURE 1.23 Schematic view of PC, directional solidi�cation, and single crystal-casting practice.




PC casting requires closed molds that may include shell, core, heat blanket, gates, and so on. DS and SX molds are designed in open-ended manner in which, the bottom of the shells end up with a water cooled chill block instead of shell to start and propagate solidi�cation in the preferred direction. DS or SX process needs additional heat source to prevent unwanted solidi�cation in mold inner walls hence mold preheat temperature is higher than that of molten metal and PC molds. As solidi�cation starts and progresses from chill block, the molds should be removed with proper velocity from the heat source to allow solidi�cation. The solidi�cation rate and speed of withdrawal must be equal. The solidi�cation rate must be balanced to a certain range so that it would be pos-sible to avoid both coarse dendritic macrosegregated metal (too slow) and secondary nucleation in front of the solid liquid interface (too fast). The main difference between DS and SX mold designs is additional selector in SX to reduce the number of solidi�ed grains to mono. As can be expected easily, fully controlled process automation is used in DS or SX casting like many other production steps of superalloys.

Casting defects may occur in production of superalloys; some of them are common to casting process thus may also be observed in other alloying systems. Porosities, core shift, core or shell breakage, inclusions, severe segregation, hot tears, un�lled metal in thin sections, and so on are some of the important common investment-casting defects. Some of the casting defects are peculiar to DS or SX process such as misoriented or equiaxed grains and grain boundary cracking.

Shrinkage porosities mainly originate from volume contraction of frozen metal (macroshrink-age), and un�lled metal in interdendritic areas during solidi�cation (microporosity). Both defects can be reduced with carefully adjusted design parameters and selection of suitable gates, feeders, mold and molten metal temperature, shell thickness, and so on. Due to macroshrinkages generally being located in latest solidi�cation area, exothermic material is placed after pouring in gates to retard solidi�cation in some applications. Microporosities can be reduced or eliminated by adjust-ing cooling gradient and solidi�cation rate during casting. Post hot isostatic pressure (HIP) treat-ment is also successively used in elimination of microporosities.

Although the primary source of inclusions is alloy cleanliness, they can also come from various sources. The impurities that result from alloy contamination can be greatly reduced by vacuum melting and double or triple melting steps. Besides, ceramic �lters placed on gates �lter off large inclusions during pouring. Even if the alloy cleanliness is ensured, unwanted chemical interactions between molten metal and crucible, shell, and core can still be an additional source of inclusions. During the DS and SX casting, the contact time between the molten metal and the shell and the core are considerably longer, thus these unwanted reactions can easily occur. In addition, steel tubes used as ingot molds may be another source of contamination if not cleaned thoroughly. Extensive studies are still ongoing today to further improve investment-casting technology.

1.4.2 WROUGHT ALLOYS

Although early superalloys have been produced by casting, many of the parts are now produced by plastic-working techniques today such as rolling, forging, and so on. Plastic working-processes do not only aim to achieve �nal shape, but also desired properties and re�ned microstructure. Alloys which contain high amount of γ ′ can be produced by plastic-working techniques, but with dif�culty. Some alloys that contain above 40% are even impossible with plastic working. These alloys can only be produced by casting or powder metallurgy techniques. The increasing amount of γ ′ means increasing solid solution and decreasing incipient melting temperatures thus the alloys hot work-ability range is signi�cantly narrowed. For example, for Inconel 718 alloys with high Ti amount, rolling operations can be performed typically between 980°C and 1040°C. Higher γ ′ also means higher strength and lower ductility. Another dif�culty in hot working superalloys is their strain hardening behavior: superalloys rapidly harden under strain due to precipitation and dislocation interactions. Trace elements such as Pb, Ag, Bi, and some intentionally added minor elements such as Zr signi�cantly reduce workability (Jackman 1984).



Generally, wrought alloys exhibit better microstructural homogeneity, whereas cast alloys show better stability under elevated temperatures. Plastic deformation is also quite useful for eliminating microporosities in the alloy.

During a hot deformation process, the alloys are subjected to recrystallization and recovery. These microscopic changes are affected by strain rate, strain ratio, residual strain from previous processing steps, and temperature. The reduction ratio should be limited between the heating steps; the heating rate should also be slow enough to prevent the alloy from cracking due to thermal stress. Three types of recrystallization (RX) occur in hot-worked superalloys. Dynamic recrystallization occurs during deformation (strain) of heated part, metadynamic recrystallization occurs just after the working due to superimposition of residual strain and thermal strain resulting from cooling (still suf�ciently hot) of the part and static recrystallization occurs without deformation.

Superalloys can preserve considerable strength at elevated temperatures thus plastic deformation becomes dif�cult. Different techniques such as superplastic forging, combination of powder metal-lurgy and forging are frequently applied to superalloys to overcome this dif�culty. Furthermore, by virtue of computer simulation on stresses, operation of the parts, designing of the molds and determination of the process parameters. . .all become more reliable.

1.4.2.1 Mill Products: Primary Hot WorkingCast ingots are converted to the semi�nished �at or round milling products such as sheet, plate, bar, billet, and so on by plastic-working operations like forging, rolling, and extrusion. While rolling and forging operations are performed in multiple steps, the extrusion process can convert an ingot to �nal shape at one go. These semi�nished milling products are transformed into the �nal parts by forging, forming or welding, brazing and machining, and so on.

Homogenization heat treatment is applied to the ingots under elevated temperatures for a long duration of time to eliminate interdendritic and grain boundary segregations. The quality of the �nal product depends on all process steps from ingot to deformation or heat treatment steps. The �nal steps are particularly critical and need to be carefully performed to obtain the best results.

The wrought and cast versions of the superalloys have usually small differences in alloying ele-ments to improve workability. Especially C may be less in wrought version to decrease hard carbides, which signi�cantly reduces workability. As superalloy properties are closely related to the grain size, all thermo-mechanical operations are performed with the desired �nal grain size in mind.

1.4.2.1.1 RollingRolling is applied to superalloys under elevated temperatures to obtain both �at products such as sheets, plates, and round products such as shaped bars. Rolling requires big forces, tough and wear, and temperature resistant rollers. The schematic view of a rolling process is given in Figure 1.24.

Sometimes the alloys being rolled are encased to improve alloy �nal surface quality. (Donachie and Donachie 2002). Furthermore, rollers should be inspected frequently to maintain required

Rollers

Plate

Rolling direction

FIGURE 1.24 Schematic view of rolling process.




product surface quality (Jackman 1984). Alloys such as iron–nickel-based precipitation-hardened Inconel 718, nickel-based solid solution-hardened Hastelloy X, and cobalt-based L-605 are widely produced by rolling to obtain semi�nished milling products which are generally found in rolled and annealed condition. Final cold work reduction ratio determines hardness. Sheet metals are usually converted to the �nal product by fabrication techniques such as forming, welding, and machining. Therefore, sheet microstructural homogeneity, thickness accuracy, surface roughness, and weld-ability are usually critical and determinant of the �nal product quality.

1.4.2.1.2 CoggingThe initial ingot breakdown usually starts with cogging that converts ingots to billets. Cogging is a special form of forging process using open dies and a large hydraulic forging press which breaks down the severely segregated dendritic as cast microstructure by successive heating and working cycles. The workpiece is heated suf�ciently to allow recrystallization. The die mate-rial is generally made from high-strength superalloy such as Waspaloy to minimize tool erosion and may be heated for some applications. The �nal shape can be achieved with multiple pass. Generally, each successive pass is at a lower temperature to obtain desired re�ned grain struc-tures. Grain re�nement gives possibilities of lower �ow stresses, more uniform reductions during subsequent forging and better ultrasonic inspectability. A schematic view of cogging process is given in Figure 1.25.

Usually round, octagon, or rectangle with rounded corners-shaped billets are produced in cog-ging. The main cogging defects are lapping, chilling and thermal cracks. Sharp corners should be avoided to prevent thermal cracks during rapid cooling. The dies must have proper radii to prevent lapping and cold die–hot workpiece contact time should be minimized to avoid chilling defect, which occurs as a result of fast cooling of the workpiece surface. Die preheating is bene�cial to pre-vent chilling in superalloys. The forging force should be high enough to cause strain at the interior of the workpiece.

1.4.2.1.3 ExtrudingExtrusion has advantages of fast one pass production of semi�nished materials with high reduc-tion ratios. Alloys to be extruded are frequently placed in a stainless or mild steel can before the extrusion process to protect tooling and improve surface quality of the alloy. Lubricants are also necessary to reduce friction and glass is widely used for this purpose. The extrusion process of

Load direction

Ingot

Hydraulic press ram

Open die

FIGURE 1.25 Schematic view of the cogging process.



superalloys needs very high capacity press and can only be applied to certain alloys that have good hot workability. Udimet 700, Astroloy, A 286, and Inconel 718 are typical examples of extruded alloys (Jackman 1984; Donachie and Donachie 2002). A schematic view of extrusion process is given in Figure 1.26. Besides solid shapes, hollow shapes such as seamless tube are also produced with extrusion. Nevertheless, the main application of extrusion process in superalloy world is pow-der consolidation.

1.4.2.2 ForgingThe intention of secondary plastic work is to give the �nal shape of the parts with desired proper-ties. Various techniques such as forging, forming, and so on have been widely used for this purpose. In working of superalloys, usually heating is involved with deformation steps. These thermome-chanical steps greatly effect the microstructure and �nal properties. Cold deformation is used only on sheet products with speci�c properties that are achievable only through cold work.

Forging is an important thermomechanical process for superalloys aiming at grain re�nement, microstructural arrangement, desired grain �ow and, of course, obtaining the desired shape. For early superalloys, forging was not too dif�cult, usually not too much different from that of stainless steels. But with introduction of precipitation-hardened alloys, forging became dif�cult and in some of today’s superalloys, such as those containing high amount of γ ′, forging can only be done via powder metallurgy techniques. Recrystallization during a forging sequence is necessary to reduce strain to improve further workability, obtain desired grain size, and eliminate grain or twin boundary carbides. The recrystallization rate is directly proportional to the temperature and the deformation amount. Grain boundaries are the preferred nucleation sites for recrystallization (Bryer et al. 1985).

Hydraulic press is the most widely used tool for forging of superalloys. With a hydraulic press, big forces can be applied to the parts in very short times. Mechanical screw-driven presses can also be used. The parts are usually heated in furnace near the forging press. The furnace tempera-ture is usually set higher than the desired temperature of the part to reduce furnace-soaking time and ensure that the interior sections of thick parts reach up to the surface temperature in short time. A note should be made that as thickness of the work parts reduces, the furnace soak time decreases and cooling rate increases during forging. The parts are heated above forging temperature to compensate for cooling as the parts immediately start to cool after removing from the furnace. Generally, the average value of starting and �nishing temperature of the part is taken as forging temperature and it is desired to keep these values minimum. The required heat is supplied through-out the forging process to precisely control the temperature in isothermal or superplastic forging by special equipment instead of a separate furnace.

The main methods for forging of superalloys are die forging, upsetting, extrusion forging, roll forging, and ring rolling. Usually parts are produced with multiple steps by combining different forging methods. For example, extrusion or upsetting can be used to make preform to parts for further closing die operation. Die forging can be performed with both open and closed die con-�gurations. Ring rolling usually produces much greater deformation than other methods and hol-low axisymmetric parts such as turbine and compressor casings can be produced by this method. A schematic view of forging steps for a turbine disk is given in Figure 1.27.

Extrusion block Extruded superalloyPiston

Load direction

Cylinder

FIGURE 1.26 Schematic view of the extrusion process.




Isothermal forging or hot die-forging technique keeps temperature constant or in tight range during forging. As the temperature of the parts to be forged can be more precisely adjusted, reliable and reproducible results are achievable with desired microstructure. Besides, as die chilling does not exist in this method, slower strain rates become possible which means more recrystallization can occur. The increasing amount of recrystallization makes forging massive parts with less energy possible. Other important advantages of the isothermal forging include the ability to form complex pro�les that cannot be obtained by conventional forging, reduction of the amount of input material, and elimination of the intermediate forging steps (Kulkarnı 1983).

The development of gas turbine design leads to producing new superalloys of higher mechanical strength and better resistance of elevated temperature. These properties mean lower forgeability. At the same time, these turbine parts also have poor machinability. Thanks to computer-aided design, the forging process can be simulated and critical information obtained before the actual forging operation (such as material �ow amount and distribution, starting and �nishing temperature dis-tribution, etc.) Very expensive and time-consuming die design is also performed more easily and accurately with computer-aided design and fabrication.

1.4.2.3 FormingAs in other metallic alloys, forming processes such as drawing, spinning, press-brake, and stretch forming all apply to superalloys that have relatively thin sections, namely, gas turbines, which exten-sively use sheet metals. Forming of superalloys is an important metal-shaping technique to give bends, curved surfaces to superalloy parts. Although, superalloys can be formed both under hot and cold conditions, cold forming is usually preferred especially in thin sheets because of the relatively narrow hot-forming temperature range (about 925–1260°C) and process dif�culties (Donachie and Donachie 2002). Depending on deformation amount and alloy properties, cold work usually includes intermediate annealing treatment steps to soften the alloy. The annealing frequency and alloy formability are greatly determined by the work-hardening rate. Low work-hardening rate alloys such as A 286 and Hastelloy X can be reduced by up to typically 90% before annealing.

Billet

Ring rolled Machined (EDMcutting, turning, milling,

and broaching)

Cogged (pancake) Forged

FIGURE 1.27 Forging steps for superalloy turbine disk.



Whereas this reduction values in Rene 41 and Udimet 500 are only 35–40% (Campbell 2006). Annealing treatment is usually performed below the solution temperature and should produce enough ductility. Annealing should not be applied to the alloys that have undergone cold deforma-tion below 8-10% to prevent abnormal grain growth. The cold work-hardening rates of various superalloys and steels are given in Figure 1.28.

Superalloys’ work hardening rates are closely related to their chemical composition. Cobalt-based alloys usually require more power than other groups of superalloys. Generally, W, Mo, Si, and C elements decrease alloy formability. Carbides greatly reduce ductility and cause cracks in cold deformation. Cobalt-based alloys usually contain much more C than nickel- and iron–nickel-based alloys. As can be expected, phase structure also affects formability, increasing volume frac-tion of γ ′ requires greater force. After the last annealing, one �nal shaping is done to give the part its exact �nal shape. The �nal cold work (usually very small) may be left to the part to avoid possible distortion during further annealing. Galling refers to the adhesive wear and transfer of material between metallic surfaces in relative converging contact during sheet metal forming. Galling is a typical surface defect in forming of superalloys. To overcome galling, tools and dies are usually plated with chrome, lower deformation speeds are selected, and lubricants are used (Campbell 2006).

0

150

200

250

300

350

Har

dnes

s (H

V)

400

450

500

550

Alloy 188

Alloy 230

Alloy 214 Alloy X

AlloyX-750

Alloy 600and

Alloy 800

A-286

Low-carbon ferritic steel

Type 304stainless steel

10010 20 30 40

Cold reduction (%)50 60 70 80

FIGURE 1.28 Cold work hardening rates for various superalloys and steels. (Reprinted with permission of Donachie, M.J., Donachie, S.J. 2002. Superalloys A Technical Guide, p. 107. ASM International: USA.)




1.4.3 POWDER METALLURGY

Powder metallurgy techniques have been widely applied to superalloys since the 1970s. The main motivation for using powder metallurgy techniques for superalloys is to �nd a way out to pro-duce hardly forgeable disks made from high-strength nickel-based superalloys. As mentioned in the investment-casting section, casting techniques are not ef�cient to obtain re�ned, chemically homogenized �nal structure in massive parts. The alloys that are poorly forgeable such as Rene 95 or Inconel 100 can be produced with powder metallurgy technique. Powder metallurgy gives possi-bility to superalloys that contain higher amount of alloying elements and greater high-temperature capability (Campbell 2006). HIP superalloy products are very close to the �nal shape of the part (near net shape). As such, deformation process, material consumption, and number of processing steps can be minimized. Powder metallurgy techniques made it possible to produce very �ne, homogeneous microstructures with less segregation and uniform phase distribution. This micro-structure greatly increases workability of the parts and allows superplastic deformation (Tien and Caul�eld 1989).

Superalloy powders are usually produced by gas atomization techniques. Argon atomization (AA) is the most widely used technique in superalloys. A schematic view of the process is given in Figure 1.29. As can be seen in this �gure, molten metal is poured into a ceramic tundish under vacuum then melting stream breaks up the molten metal into droplets via high-pressure argon gas jets, �nally the molten metal droplets solidify with cooling rates about 100 K/s in the atomization chamber.

The typical appearance of gas-atomized powder particles is shown in Figure 1.30. The attached smaller particles are called satellite, which result from collision between solidi�ed powder and par-tially molten droplets can also be seen in this �gure.

Vacuum environment

Induction furnace

Molten metal

Tundish

Solidified metal droplets

Atomization chamberExhaust

SecondarycollectorPrimary

collector

Cyclone seperator

Gas jet

FIGURE 1.29 Schematic view of the argon atomization process.



The other important superalloy atomization process is vacuum atomization (VA). In this tech-nique, both melting and powder collection tank work under vacuum at the beginning of the process. The alloy is inductively melted under vacuum then the melting chamber is pressurized with hydro-gen. Once pressurization is complete, a sealing valve that connects two chambers is opened and molten metal is sprayed into the collecting tank and is solidi�ed there.

The main centrifugal atomization techniques are rotating disk, rapid solidi�cation rate atomi-zation, and rotating electrode. In rotating disk method, vacuum induction-melted superalloys are poured into a tundish and metal stream falls onto a rotating-chilled disk. Molten metal breaks up into small droplets, which are accelerated to the sidewalls of inert gas-cooled chamber by centrifu-gal force. A schematic view of the rotating disk process (RDP) is given in Figure 1.31.

A schematic view of the rotating electrode process (REP) is given in Figure 1.32. In this method, an arc is created between consumable rotated electrode and tungsten tip. The molten metal droplets are dispersed by centrifugal forces and solidify in �ight in atomization chamber. A plasma arc can also be used instead of a tungsten arc to melt the consumable electrode. The cooling rate is signi�-cantly higher than gas atomization with a value of about 105 K/s. The early production of IN-100 and Rene 95 powder was performed by rotating electrode technique. Generally, REP is not used to produce superalloy powder today (Donachie and Donachie 2002).

Alloy and process cleanliness is extremely important in powder metallurgy. Organic or inor-ganic contaminations act as a source of defects like prior particle boundaries (PPBs). Entrapped gas in powder particles (hollow particles) can cause porosities. Ceramic and metallic inclusions may be found in powder. A variety of control tests are applied to the powder to ensure the quality of the product. The produced powder is usually stored in controlled or inert gas environment (Sims et al. 1987).

The MA process is neither powder production nor consolidation. The aim of the process is to obtain alloyed powder containing �nely dispersed oxide particles as a strengthening agent. In this method, prepared powders by various techniques such as gas atomization are blended to the desired compositions with oxide particles and charged into a ball mill. The initial powder can be found either as individual element or in prealloyed form. The ball mill containing steel balls and a rotat-ing rod with a number of branched arms is called “attractor.” During milling, the steel balls smash each other with powder particles in between, or powder particles hit each other, weld and fracture

FIGURE 1.30 Typical appearance of gas-atomized Inconel 718 powder particles.




repeatedly throughout the process. Chemical composition and phase distribution of the powder particles become homogeneous and oxide dispersoids are �ner as the milling process continues. Once the desired properties are achieved, mechanically alloyed powder is removed from the ball mill then standard consolidation techniques are applied to produce the parts. A schematic view of the ball mill is given in Figure 1.33.

Tundish

Molten metalstream

Y

X

Powder droplets

Induction furnace

Vacuumenvironment

Vacuum pump

Molten metal

Rotating disk

FIGURE 1.31 Schematic view of the rotating disk process (RDP).

Tungstenelectrode (–)

Vacuum pump

Powder droplets

Rotating mechanism

Powder collection tank

Rotating consumableelectrode (+)

Vacuum environments

FIGURE 1.32 Schematic view of the rotating electrode process (REP).



1.4.3.1 Powder Consolidation TechniquesThe superalloy powders produced are converted into billets by extrusion or HIP process. These billets are further processed thermomechanically by isostatic or hot die forging to obtain �nished parts. The parts can also be obtained directly by near net shape HIP process without forging. Superalloy powders are also consolidated by various ways such as pressing and sintering, forg-ing, vacuum hot pressed, and so on. Regardless of powder consolidation route, sometimes �nal machining is required to give precise dimensions to the �nished part. The aim of the consolida-tion process is not only to give �nal shape to the part, but also to obtain desired properties and microstructure. The PPB and porosities should be eliminated with consolidation process. The main causes of PPBs are powder contaminants, which collect on the powder surfaces and act as carbide nucleation sites. Therefore, the C content of the P/M alloys are signi�cantly lower than cast counterpart.

1.4.3.1.1 ExtrusionPowder can be extruded directly or hot compacted before extrusion process. Direct extrusion is per-formed under high temperature with a reduction ratio of about 13:1. In another case, the powder is usually placed in a container and hot compacted then extruded to obtain billet. Hot compaction pro-cess usually performs under solution temperature and approximately 95% density can be obtained. A fully dense billet can be obtained with a typical extrusion ratio of about 6:1.

Extrusion is also used for consolidation of mechanically alloyed ODS powders. ODS powders have nonreactive properties due to their oxide nature. The powder can be cold compacted by cold isostatic pressing (CIP) or uniaxial pressing. Loose powder can also be used directly for extrusion process. After the powder is canned and heated, extrusion is applied to obtain full density. Extrusion rate, reduction ratio, and extrusion temperature are the key parameters of the extrusion process that greatly affect the �nal properties of the product.

1.4.3.1.2 Hot Isostatic PressIn HIP process, powder is packed in a sheet metal container and placed in HIP chamber. Containerized powder is usually evacuated before sealing. Powder is pressurized by gas (usually argon or helium) and heated in HIP chamber. The typical HIP process pressure value is 15 ksi and temperature value is about 1100–1200°C. HIP temperature can be adjusted either above or under γ ′

Controlledatmosphere

Rotating rod

Attractors

Steelballs

FIGURE 1.33 Schematic view of a ball mill.




solution temperature depending on the desired microstructure. Temperature, time, and pressure are the key parameters of the HIP process.

1.4.3.1.3 Thermomechanical ProcessingBesides shaping, thermomechanical process signi�cantly improves alloy properties such as fatigue strength by eliminating or reducing contamination-related defects. The defects are broken up by metal �ow and dynamic recrystallization during thermomechanical processing (Sims et al. 1987). In isothermal forging, both die and alloy have the same temperature while die temperature is lower in hot die forging. The most common die material in thermomechanical process is TZM molybdenum with excellent high-temperature properties (Campbell 2006). Unfortunately, vacuum is needed to protect these properties thus treatment cost is high. The required force is signi�cantly lower than conventional forging due to superplastic behavior, which makes near net shape produc-tion possible.

The �nal properties of the P/M superalloys depend on alloy composition, powder particle size, consolidation processing, the degree of defect free structure, and heat treatment. P/M superalloy’s microstructure has more uniform and �ner grain size than wrought and cast alloys. Owing to �ner grain size, P/M superalloys exhibit better ductility and tensile strength.

1.4.4 WELDING

As mentioned before, the most common applications of superalloys are gas turbine engines thus weight is an important consideration especially in aviation gas turbines. Therefore, the designers must rely on superalloy components’ properties to adequately approach their design limits. Usually gas turbine components include joint sections, thus the design requirements should also be met in these sections. As other engineering alloys, superalloys can be joined in various ways such as weld-ing, brazing, fastening, and so on. Apply the principle of “the weakest link of the chain” to the gas turbine parts and the critical importance of joints is clearly understood.

As in many other processes, γ ′ amount has critical importance on weldability of precipitation-hardened nickel- and iron–nickel-based superalloys. Solid solution superalloys are easily weldable without any special pre or post-heat treatment and usually can be used in as-welded condition; whereas precipitation-hardened superalloys are usually welded in solution annealed or over aged condition.

The increasing volume fractions of γ ′ make welding dif�cult. Two types of cracks may occur in welding of superalloys. These are hot cracking, seen during welding and post-weld heat treatment (PWHT) or strain age cracking, usually observed after weld heat treatment as its name implies. As the main γ ′ forming elements of superalloys are Al and Ti, several studies investigate the effect of these elements on weldability. The effect of aluminum versus titanium content on weldability for superalloys is illustrated in Figure 1.34.

When Figure 1.34 is examined, alloys such as IN-100 and Inconel 713C with high titanium and/or aluminum contents are considered unweldable. Generally wrought alloys that have more than 0.35 volume fraction of γ ′ are sensitive to hot cracking. Hot cracking usually occurs during welding in the heat affected zone (HAZ). Hot cracking can also be seen in many alloy types such as steels and usually stems from a combination of high thermal strain and local melting of low-melting point phases (frequently in grain boundaries) in HAZ. Besides unwanted trace elements, carbides and borides in the grain boundaries are preferred local incipient melting areas of superalloys in HAZ. Proper and careful design of weldments, using high-purity alloys considerably reduces hot cracking in superalloys (Owczarski 1984).

The main reason for poor welding properties of γ ′ hides behind aging phenomena. The welded alloy should frequently be subjected to solution and aging heat treatment to reorder the micro-structure in weld area and relieve residual stress after welding. High γ ′ volume fraction alloys have less strain tolerance and high residual stress at welding areas due to fast thermal cycle and



metallurgical phase changes. The alloys have to pass the aging zone to reach solution temperature and due to very fast aging characteristic nature of γ ′, additional stress is caused within the alloy and cracks occur in base metals that have high volume fractions of γ ′. Usually, alloys that are sensitive to strain-age cracking are over aged before welding to overcome this problem at least to some degree. Unfortunately, over aging tactic is not useful in HAZ because the thermal history, which comes from welding process, rearranges the microstructure in this area. Unlike γ ′, γ ″ is not subjected to strain age cracking, thanks to the slower precipitation rate (Campbell 2006). This is another important issue and one of the reasons why Inconel 718 alloy is so popular among the precipitation-hardened superalloys. Inconel 718 is readily weldable. Before solution and after stress relieving plus aging, heat treatments sequence give best result for Inconel 718 welding practice. As for other welded materials, proper cleaning before welding is critical in superalloys. The alloys should be completely free from oxides and dirt. As superalloy oxide �lm strongly adheres to the surface, powerful cleaning methods such as abrasive grit blasting instead of simple wire brushing are recommended. Any oxide residue may cause insuf�cient fusion in weld/base metal interface and entrapped oxides in weld metal. The alloys are also readily oxidized during welding and should be protected by an inert or oxidation resistant environment. The most widely used protective gases for this purpose are argon, helium, and nitrogen. Although helium is the most expensive among them, it has advantages of reduced porosity and faster welding. Argon gives better results than nitrogen but is of higher cost. In case of insuf�cient protection, oxide �lm may form on the weld surface or oxide �akes in the weld metal. Superalloys are usually welded by multipass application to minimize heat

Inconel 713C

8

7

6

5

4

Ti %

(wt)

3

2

1

0 1 2 3Al % (wt)

4 5 6 7

Inconel 100

Inconel 738

Inconel 702

Inconel 718

Inconel X-750Inconel X

Hastelloy X

Readily weldable

Rene 95

Rene 41

WaspaloyNimonic 80A

Difficultly weldable due tostrain age cracking

Mar-M246

Mar-M200Udimet 600

Udimet 700

Udimet 500

Unitemp 1753

Astroloy

FIGURE 1.34 Effect of aluminum versus titanium content on weldability for superalloys. (Adapted from Thompson, R.G. 1993. ASM Handbook Vol. 6 Welding Brazing and Soldering: Welding Metallurgy of Nonferrous High Temperature Materials. ASM International; Campbell, F.C. 2006. Manufacturing Technology for Aerospace Structural Materials. Elsevier; Donachie, M.J., Donachie S.J., 2002. Superalloys A Technical Guide. ASM International: USA.)




input and residual stress, therefore oxide �lm, which forms during welding, is entrapped in passing borders and acts as stress raisers together with oxide �akes.

1.4.4.1 Gas Tungsten Arc Welding and Gas Metal Arc WeldingThe most widely used welding methods in superalloys are GTAW (gas tungsten arc welding), also referred to as TIG (tungsten inert gas) welding and GMAW (gas metal arc welding). A schematic view of the GTAW process can be seen in Figure 1.35. In this technique, an arc is created between workpiece and nonconsumable tungsten electrode by using direct current. Tungsten electrode mate-rial is usually alloyed with 1–2% thorium, cerium oxide, or lanthanum oxide. The base metal is melted and �ller metal is fed in this arc manually or automatically. Some applications do not need �ller metal especially during welding of thin sections.

Protective gas is supplied by weld torch and covers the weld area. Beside nozzle shield, which protects weld bead and HAZ, trailing and back shields are usually used to protect hot solidi�ed weld metal and root of the weld, respectively. Flow rate of the protective gas should be properly selected. Too low or too high a �ow rate causes insuf�cient protection due to ineffective purge or air in�ltration and gas turbulence, respectively. As seen in Figure 1.35, the workpiece is hotter than the electrode due to the workpiece being positively charged. Therefore, GTAW process can produce desired narrow welds with deep penetration, which means minimum HAZ and lower heat input to the base metals.

In GMAW process, the �ller metal is directly used as electrode instead of tungsten electrode. Another important difference is the electrode being positively charged (anode) unlike GTAW. As positive electrode is hotter than negative workpiece, consumable �ller metal completely melts eas-ily and quickly. Consumable wire electrode feed rate is usually automatically adjusted by weld-ing machine with feedback control system. The main advantages of GMAW process over GTAW are energy ef�ciency (melted metal amount per unit power consumption) and higher welding rate. The ef�cient melting of �ller wire is especially suitable for welding of thick plates (above 10 mm) (Campbell 2006).

In case of welding of superalloys with similar chemical composition, �ller metal composition is also usually the same with a small addition of deoxidizing elements, which help to protect weld quality. The welding of dissimilar superalloys is a more complex phenomena and the �nal chemi-cal composition of weld metal is a mixture of two base metals and �ller material. The dilution amount of base metals may show difference, furthermore other factors such as joint design and welding techniques also affect weld composition. The �nal properties of this composition should have proper chemical, mechanical, and physical properties. For example, thermal expansion coef-�cient difference between base metals and weld metal causes additional stress in high-temperature service environment and especially reduces fatigue life. Nickel–chromium-based �ller electrodes

Sheet thickness

Penetration

TorchTunsten electrode [(–) cathode]

Welded material [(+) anode]Arc

Leg size

Concavity

Gas shield

�roat s

ize

FIGURE 1.35 Schematic view of the GTAW process.



are widely used in dissimilar nickel alloys welding. Joint design is also an important tool to reduce unwanted stress in the weld area. Postweld cleaning process is also important as in cleaning of base metals before welding. The cross section of real �llet TIG-welded microstructure of Inconel 718 can be seen in Figure 1.36.

1.4.4.2 Resistance WeldingAnother fusion welding process is resistance welding. As resistance welding is well suited for sheet metals and superalloy gas turbine components are generally also made from sheet metals, this tech-nique is widely used. The main principle of the application is based on a combination of pressure and electric current to the overlapped parts that are going to be welded. The electrical resistance between mating surfaces of the parts creates heat and causes local melting. The solidi�ed weld area is called “nugget” and usually does not reach to the surfaces of the parts. The nugget size and shape depend on the electrode shape and current amount. The resistance welding process can be used to produce weld nuggets of both spot and seam form, which depend on continuous or discrete electrode action and movement of the parts. A schematic view of the resistance welding process and resulted microstructure can be seen in Figure 1.37.

As superalloy components are expensive and usually used in critical applications, more sophis-ticated welding methods can be used to obtain better joints. Electron beam welding (EBW), laser welding, inertia welding, and friction welding are some of the used application examples of different welding techniques of superalloys.

1.4.4.3 Electron Beam WeldingIn EBW, an electron beam under vacuum is used as a heat source. The focused beam has a very high energy that easily melts and even vaporizes base metal. The schematic view of the EBW machine is given in Figure 1.38. Filler metal normally is not used in this technique. The main advantage of EBW process is the capability of very narrow weld bead with huge penetration values thanks to extremely thin and energized nature of the electron beam. This welding shape indicates that it is pos-sible to join thick sections with minimum heat input. The �anges of the turbine nozzles, cases, and disks can be perfectly welded by EBW technique. The more conventional methods such as GTAW are not capable of this type of weldment without damage to the component. A schematic view of the turbine nozzle with an EBW-welded �ange section is given in Figure 1.39. The microstructure of the electron beam-welded superalloy can be seen in Figure 1.40. The main disadvantages of the EBW process are the high initial and maintenance cost of the machine and manipulating dif�culties of the workpiece due to vacuum environment and precise �xture requirement.

FIGURE 1.36 Cross section of �llet weld microstructure of Inconel 718 (15× and 50×).




1.4.4.4 Laser WeldingIn laser welding, a laser beam is used as a heat source instead of electron beam. Nd-YAG and CO2 lasers are widely used for this purpose. The other aspects of the process are very similar to EBW, but minus the vacuum chamber. The initial and maintenance cost is also high.

1.4.4.5 Friction WeldingAn interesting application of friction welding in superalloys is the production of integrated one piece engine rotors (blisk). The blades are attached to the disk by this method. The combination of a linear oscillary motion and compressive force produce heat and solid state joint. Inertia welding is very similar to friction welding where rotational symmetric parts are bonded without melting (Campbell 2006). Precipitation-hardened nickel alloys such as Waspaloy, Astroloy, and so on can be joined by

FIGURE 1.37 Schematic view of the resistance welding process and resulted microstructure.

Tungsten filamentAnode

Observation window

Electron beam

Focusing coil

Deflection coil

EB-welded workpiece

Vacuum environment

5 axis table

FIGURE 1.38 Schematic view of the EBW machine.



inertia welding. A good example of inertia welding process is a number of gas turbine disks that are bonded together to obtain signi�cant weight saving with inertia welding. The main disadvantages of the method are very high initial cost of the plant and limited number of applications.

1.4.4.6 ConclusionWelding is a critical process that changes the properties of the superalloys substantially; therefore, all of the welding process steps should be carefully designed. For instance, the blending of weld reinforcement in butt welds usually gives better fatigue properties due to reducing geometric stress concentration area. The static properties of alloys such as tensile strength are not usually affected by blending, while the creep property of the weld is usually reduced due to reduced section area. Therefore, the requirements and the result of the processing step of the weldment should be well understood. The initial strength of many superalloys can be nearly obtained by proper welding and heat treatment processes.

1.4.5 BRAZING

Brazing process is very similar to soldering but with higher bond strength and application tempera-ture (above 450°C). The application temperature should be under base material incipient melting point and enough to melt braze material. Usually nickel or cobalt-based brazing alloys are used for brazing superalloys and some elements such as Si, B, and P are added to the braze material to reduce

FIGURE 1.39 Schematic view of the turbine nozzle with an EBW welded �ange section.

FIGURE 1.40 Microstructure of the electron beam welded superalloy (50×).




the melting point. Proper brazing materials depend on both base metal chemistry and application needs. Besides their base metals and melting point depressants, braze materials usually contain Cr to improve corrosion resistance. Braze materials should have some speci�c properties such as no detrimental effect on base metals alloy chemistry, high wetting capability, proper melting point, good mechanical and chemical resistance, good corrosion and creep resistance, and so on.

A schematic view of the brazing process application steps is given in Figure 1.41. The brazing appli-cation includes cleaning of the parts, �xing of parts and applying of braze materials in the joint area, applying heat, cooling down the parts to be brazed, and blending of the reinforced or drained area.

Cleaning is very important in the brazing process and directly affects the bond quality. Incomplete fusion and porosity defects are strictly related with precleanliness of the parts and process com-ponents. Besides mechanical and chemical cleaning actions, vacuum cleaning is widely used in cleaning of superalloys and is a very effective way in eliminating surface oxide �lms and dirt of the hardly reached areas such as inner surface of the cracks.

A clearance remains between parts to be bonded in the joint area and brazing material is put into this clearance, then heat is applied to melt braze material. Sometimes it may also bene�t from capil-lary effect for �lling this clearance especially during repair of cracks. Both chemical and physical bonding can be obtained via brazing. Braze material can be found in various forms such as wire, sheet, folio, powder, and paste. Powder form is usually mixed with binder to obtain slurry. The binder should not leave residue at the end of the process. Acrylic-based materials are widely used as binder.

After application of braze materials to the joint areas, the parts are furnaced. Superalloys are usually furnaced under vacuum or inert atmosphere. Although the furnace environment is protec-tive, the surface of the precipitation-hardened alloys, which contain highly reactive elements such

Chemical and mechanical cleaning(acid or alkali dipping, grit blasting, wire brushing, etc.)

Furnace cleaning(under vacuum or inert atmosphere)

Preparing braze mixing(powder + slurry)

Applying braze material to the jointarea (injecting, brushing, etc.)

Applying stop offmaterial (if necessary)

Drying (usuallyunder air)

Furnacing

Grinding excessive braze(if necessary)

Postcleaning (remove excessiveresidues, stop off material)

Inspection (evaluate braze area against defects such asvoids by radiographically and/or metallographically)

FIGURE 1.41 Schematic view of the brazing process application steps.



as Al and Ti, is slightly oxidized. This oxide layer can adversely effect the wetting capability of the braze material, therefore usually �ux material is added to the braze �ller metal to improve wetting. Cobalt-based superalloys can also be brazed under reducing atmospheres such as hydrogen.

Brazing made possible the joining of precipitation-hardened superalloys, which contain high amount of γ ′ otherwise not able to weld. The main disadvantages of brazing compared to welding are brittleness and loss of strength in high temperatures due to low-melting point of the braze mate-rial. Application areas vary: turbine vanes airfoil—outer and inner band connections, turbine cases, vane segments, and diffusers joint sections, the repairing of turbine blades and vanes airfoil sections, and so on. The brazed airfoil/outer band connection of turbine vane can be seen in Figure 1.42.

Another important solid state joining technique is transient liquid phase (TLP) bonding, also known as diffusion brazing. While brazing is diffused only in limited amount into the base metal, TLP is completely mix �ller metal and base metal with the result of long isothermal heat treatment. Although �ller material melting temperature is low like that of braze material in the beginning of the process, the joint melting temperature signi�cantly increases after TLP bonding as a result of diffusion process. The treatment time can vary signi�cantly and can be between 1/2 h and 80 h or even longer. Main process parameters are furnace time, application temperature, mutual solubility of the base and �ller metals and the amount of �ller metal. The process is usually performed in vacuum furnace and after application the microstructure in the joint area is usually undistinguish-able from base metals due to diffusion and grain growth of the base metals (Campbell 2006). The main advantage of the process is the physical and mechanical properties of the joint section being very close to that of the base metals unlike brazing.

Diffusion bonding is also very similar to TLP bonding. The difference comes from absence of the �ller metal. The base metals that are to be bonded are held under high temperature and compres-sion force to obtain diffusion.

1.4.6 MACHINING

1.4.6.1 Conventional MachiningUnfortunately, the machining of superalloys is very dif�cult. The reasons behind this dif�culty are related to the superalloy characteristics such as hard intermetallic microconstituents, strong work hardening potential, low thermal conductivity (that causes heat accumulation around the cutting tool), and the ability of retaining high strength at high temperatures: all the good points about super-alloys now work against machining. Many engineering alloys soften at the tool/workpiece contact area as the contact temperature increases, which facilitates cutting. But superalloys are designed to maintain high strength at high temperatures. . . therefore, in practice, cutting tool is the one that

FIGURE 1.42 Brazed airfoil/outer band connection of turbine vane.




wears out fast. Among all the aerospace metals, only machining of Ti alloys is more dif�cult than superalloys (Campbell 2006).

Though dif�cult, all of the conventional machining methods apply to superalloys such as turn-ing milling, grinding, broaching, drilling, planning, and so on. More sophisticated methods such as water jet, laser, electrodischarge machining (EDM) and electrochemical grinding (ECG), and so on are used to overcome machining dif�culties and obtain complex geometries, which otherwise cannot be achieved by conventional methods.

Nickel- and cobalt-based superalloys are more dif�cult to machine than iron-based superalloys. Some of the iron-based solid solution superalloy machining characteristics are quite similar to that of the stainless steels. For precipitation-hardened superalloys, the main approach to machining is: use the solid solution heat treatment to soften the alloy, machine near �nal shape, age then �n-ish machining where possible. With this strategy good surface quality, minimum tool wear and minimum part distortion can be obtained. To obtain the best results, sharp cutting tools, positive tool rake angles, low machining speeds, rigid tools, and part set-ups usually improve and facilitate conventional machinability. Cutting tool materials can vary but carbide tools have better tool life and allow higher speeds than high-speed steels. The workpiece surface �nish is also better with carbide tools. The main disadvantage of the carbide tools is their brittleness, which does not allow interrupted cutting or low rigidity of the parts, set-ups or machine. CBN (cubic boron nitride) and ceramic (Al2O3, SiAlON, etc.) tooling is also possible for high rigidity and uninterrupted machining. Tool geometry is also important: sharp and deep tool marks should be avoided, which cause stress concentrations on the �nal part surface. Especially high-strength superalloys have some degree of notch sensitivity and their service conditions may cause initiation of fatigue cracks on the surface. The machining process causes a signi�cant amount of heat generation at the tool/workpiece surface, therefore surface coolants are used to protect the part and tool from excessive heat. In the event of insuf�cient cooling, the surfaces of age hardenable superalloys can be overaged or numerous shal-low surface crack network called “heat checking” can form. Cooling agents also serve as lubricants and reduce friction between the part and the tool. Cutting tools are usually coated with hard thin �lms such as TiN, TiAlN, TiAlCN, and so on to improve surface quality and tool life. Thin �lms are usually produced by physical vapor deposition (PVD) or chemical vapor deposition techniques. Both high-speed steel and carbide tools can be coated with these techniques. Al2O3 is widely used as wheel material for grinding of superalloys.

Machining process is essential for superalloys, and is widely used in many applications such as production of fastening components (bolts, �ttings, etc.), and gives �nal shape and surface quality to both investment casted and wrought parts such as turbine disks, combustor chambers, turbine casings, and so on.

1.4.6.2 Water JetWater jet technique has been used for a long time to cut and shape low-strength materials such as printed circuit boards and plastics. The addition of abrasives to the system gives opportunity to the machining of higher-strength materials, including superalloys. In this application, high-pressure �uids, which contain abrasives, are directly sprayed to the superalloy surface via a nozzle with high pressure (Johnston 1993).

1.4.6.3 Electrodischarge MachiningEDM process uses sparks to remove material. The workpiece and EDM electrode are sunk in a dielectric �uid, a small clearance is maintained, and an electric arc is created in between. As work-piece material is removed by melting, the electrode is moved to keep the clearance between elec-trode and workpiece constant. The electrode is usually made of copper or graphite. The EDM electrode gives its own negative shape to the machined material, therefore a cutting operation is usually performed by sheet and wire, a drilling operation is performed to form electrodes in the form of wire or bar. Special shapes like airfoil pro�le can also be obtained by EDM electrode.



Oil is widely used as a dielectric �uid. Unlike conventional machining, the ef�ciency of the EDM process does not depend on the hardness or heat treatment condition of the material. The only requirements are that the material must be electrically conductive and the machining area has to be accessible for the electrode. Therefore, superalloys that are hard to machine like precipitation-hardened nickel-based group are perfectly suitable for EDM machining. The process is quite fast and easy. The surface appearance and properties of the EDM-machined parts are signi�cantly dif-ferent from classical machining, due to very localized melting action. Thus, EDM machining leaves behind melted and resolidi�ed layer on the surface of the part, called “recast layer.” The quality and the acceptability of EDM operation is determined by the properties of this surface such as the recast layer thickness, the recast layer crack amount, size, and whether these cracks extend into unmelted base metal and the heat-affected zone thickness. Although it depends on the process parameters and material, the recast layer thickness is usually well below 0.125 mm and about 0.05 mm in superal-loys. The recast layer usually contains cracks due to fast solidi�cation. The surface appearance and microstructure of the EDM-machined precipitation-hardened nickel-based Rene 77 superalloy materials are shown in Figures 1.43 and 1.44, respectively.

1.4.6.4 Electrochemical MachiningThe electrochemical machining (ECM) process is widely used in machining of gas turbine parts. In ECM, an electrolytic cell is set up, which includes the workpiece, the tool, and the electrolyte. An

FIGURE 1.43 Surface appearance of the EDM-machined gear (15×).

FIGURE 1.44 Microstructure of the EDM-machined Haynes 188 material (200×).




electric current passes through this cell, which causes anodic dissolution of the workpiece. Aqueous solutions of inorganic salts such as sodium chloride, potassium chloride, sodium nitrate, and strong acids can be used as the electrolyte.

A special form of ECM is the ECG, which uses an electrically conductive grinding wheel instead of a tool-shaped contour. ECG combines the electrochemical process and grinding operation. The workpiece and tool are conducted electrically by an electrolyte, which contains positive and nega-tive ions. The workpiece is positively charged (anode) and the tool is negatively charged. The nega-tive ions transfer to the workpiece surface, causing formation of oxide �lms on the anodic surface and machining action easily removes the oxide. ECG is used very effectively in grinding of gas tur-bines superalloy honeycomb seals. Honeycomb seals are usually made from thin superalloy sheets such as Hastelloy X and conventional grinding action causes excessive burring, which causes the honeycomb to become unusable.

1.4.6.5 Laser MachiningAs mentioned in the welding section, a laser beam has high energy; besides welding, it can also perform machining operations such as cutting, drilling, and so on. Carbon dioxide and Nd:YAG (neodymium-doped yttrium aluminum garnet) lasers are the most widely used laser types in machin-ing of metals. While CO2 laser has much higher energy capability, translating to faster machining, Nd:YAG laser has an advantage of percussion drilling and cutting of metals at angles and thicknesses not possible with CO2 laser. Figures 1.45 and 1.46 show a schematic view of the laser process and a microscopic view of a hole drilled in Inconel 718 (precipitation-hardened nickel-based superalloy) with Nd-YAG laser. High-speed machining results in very rough surface as given in Figure 1.47.

1.4.7 HEAT TREATMENT OF SUPERALLOYS

Generally, three types of heat treatment apply to superalloys, that is, homogenizing (or solutioning), aging, and stress relieving. The early cast alloys were produced at aged condition as the thick mold walls effectively slowed down cooling. These alloys were used as cast. However, the properties vary even

Electric line

Cooling water Laser source Lens

Mirror

Laser beam

Gas shield inlet

Nozzle

Workpiece

Lens

FIGURE 1.45 Schematic view of the laser process.



from the same lot because there is no control in cooling rate. In those years, superalloy microstructures were not fully understood thus the effect of cooling rate was not known (Tien and Caul�eld 1989).

Some PC cast superalloys, especially cobalt-based, were used as cast without any additional heat treatment, and still are even today when a proper cooling rate can be obtained. Due to the absence of γ ′ phase, heat treatment for cobalt-based superalloys aims at arrangement of carbide structure and homogenization. Usually cobalt-based superalloys are aged at about 760°C without solution treatment to obtain discrete Cr23C6.

Solution treatment aims at obtaining over saturated solid solution by solutioning of phases in the microstructure. To achieve this, the alloy is brought beyond phase solubility limit and cooled down rapidly. Thereafter, desired size and amount of precipitates are attainable in the aging treatment that follows. Homogenization of microstructure is also achieved with solid solution treatment in cast structures where composition/phase segregation occurs. DS and SX superalloys all undergo solid solution treatment. The homogenization treatment is usually also applied to many alloys before welding to improve weldability.

The solubility of γ ′ phase varies with the chemical composition. In some alloys, complete solu-tioning of the γ ′ phase is not possible. In this case, partial solutioning treatment is done at tempera-tures well below its incipient melting point.

FIGURE 1.46 Laser-drilled precipitation-hardened nickel-based Inconel 718 superalloy hole microstructure (200×).

FIGURE 1.47 Rough surface appearance of Rene 41 (15×) cut with Nd-YAG laser at high speed.




The heating rate becomes critical when the treatment temperature is too close to the incipient melt-ing point of the alloy. The segregations that cast structures naturally have should be homogenized before maximum temperature is reached to prevent incipient melting of these low-melting point areas. Typically, solution treatment temperature may be between 1050°C and 1220°C with 2–6 h soaking. For DS superalloys, the solid solution treatment temperature is higher than PC and may even rise up to 1300°C in SX superalloys. This is possible because of the lower amount or complete absence of grain boundary elements that reduce incipient melting point in DS and SX alloys (VerSnyder 1982). As such it is possible to obtain full or high level of solutioning of γ ′ phase, which, in turn, gives rise to complete or high-level control of size and distribution of γ ′ precipitation upon aging.

γ ′ phase precipitates immediately after the cooling starts in solution treatment. Hence, a high cooling rate is required to avoid oversized γ ′. Fan circulation of inert gas is widely used to achieve proper cooling rate. Generally, cobalt-based cast alloys do not need vacuum or controlled atmosphere furnaces during heat treatment. On the contrary, however, nickel-based alloys solution treatment should always be performed under vacuum or controlled atmosphere. The usage of inert gases instead of vac-uum is more ef�cient in preventing elemental depletion particularly aluminum. The typical solid solu-tion heat treatment temperature and time applied to the various superalloys are tabulated in Table 1.9.

Stabilization heat treatment is performed at temperatures between solution and aging tempera-tures. Coarse MC carbides are transformed to �ne grain boundary carbides and γ ′ size and mor-phology are arranged at this treatment. Vacuum or controlled atmosphere is used similar to solution treatment in nickel- and iron–nickel-based alloys. Stabilization treatment improves creep and stress rupture strength. In iron–nickel-based alloys, weldability is also improved by dissolving of δ and γ ″ phases. In addition, stabilization treatment also serves as stress reliever.

The required size and amount of phases especially γ ′ and γ ″ are obtained by virtue of aging treatment. Although aging temperature is lower than stabilization, alloys are frequently aged under vacuum or protective atmosphere. Heating and cooling rates are less critical due to lower operation temperature. The typical values of aging steps of heat treatment time and temperature for various superalloys can be seen in Table 1.7.

TABLE 1.9Typical Values of Various Superalloys Solution Time and Temperature

Alloy Solution Temperature (°C) Solution Time (h)

A-286 980 1

Hastelloy X 1175 1

Haynes 188 1175 1/2

Incoloy 907 980 1

Incoloy 909 980 1

Inconel 625 1150 2

Inconel 718 980 1

Inconel X-750 1150 2

L-605 1230 1

N-155 1175 1

Nimonic 90 1080 8

Rene 41 1065 1/2

Udimet 700 1175 4

Source: Adapted from Donachie M.J., Donachie S.J. 2002. Superalloys A Technical Guide. ASM International DeAntonio, D.A. 1991. ASM Handbook Vol. 4 Heat Treating: Heat Treating of Superalloys. ASM International: USA.



Unlike precipitation-hardened alloys, solid solution alloys do not need aging heat treatment as there is no second phase to precipitate. Therefore, solid solution alloys are used as cast. However, for these alloys, solid solution and stress relieving treatments are still needed to obtain homogeneous microstructure and to reduce residual stress (DeAntonio 1991).

1.5 MAIN FAILURE MECHANISMS OF SUPERALLOYS

Failure of superalloys components always draws great attention due to their critical application areas such as aerospace, nuclear and fossil energy, petrochemical industry, and so on.

As mentioned in the �rst section of the chapter, the thing that makes superalloys “super” is their ability of retaining properties at elevated temperatures and in harsh environments. Although the resistance is “superb” and indeed very impressive, unfortunately any material, including superal-loys, cannot be completely immune to service conditions. A superalloy’s main failure mechanisms are strictly related to the service conditions, meaning time, temperature, stress, and corrosive environment. These mechanisms can affect superalloys individually or a combination of a few of them and may cause component damage or catastrophic failure of the systems under some cir-cumstances. For instance, creep and fatigue can damage components simultaneously; fatigue life is greatly affected by environmental degradation thus under corrosive environment, “corrosion fatigue” may occur.

1.5.1 HIGH-TEMPERATURE OXIDATION

Despite their resistance, superalloys are oxidized in gas turbine service conditions, especially the parts exposed to the hot gas stream. Oxygen diffuses into the metal from the surface and forms metal oxides. These oxides have different speci�c volume and thermal contraction, expansion coef-�cients, therefore they easily crack and drop, and new surfaces are oxidized again and material properties are degraded (Kircher 1989). The rate of oxygen dissolution to the base metal is greatly affected by the properties of the surface oxide �lm. Ideally, surface oxide �lm should be continuous, have low diffusivity, strongly adhere to the unoxidized metal surface, be strain tolerant, be thermo-dynamically stable, and have slow growth properties (Diltemiz 2010). Alumina (Al2O3), chromium oxide (Cr2O3), and silicon dioxide (SiO2) have the ability to form desired protective oxide scales. However, SiO2 is not suitable in superalloys due to formation of an undesirable phase detrimental to mechanical properties (Diltemiz 2011).

The selective oxidation of aluminum and chromium elements in superalloys forms protective alumina (Al2O3) and chromium oxide (Cr2O3) scale. The Superalloy’s oxidation protection is based on selective oxidation of these elements. Yttrium may improve the adherence of the protective oxide scale, tungsten, however, may cause decrease in oxidation resistance.

Chromium oxide is also very effective in hot corrosion (sul�dation) protection. As chromium oxide becomes volatile above about 1000°C, Al2O3 is the main protective oxide in superalloys, which operate at high temperatures. As oxidation continues, the concentration of protective alloying elements near the surface decreases and less protective oxides such as NiO start to form. The oxida-tion resistance thus signi�cantly depends on the initial concentration and consumption rate of these elements. Superalloys are frequently coated for maximum protection. These protective coatings contain high concentration of aluminum and chromium beside Ni or Co and may also contain small quantities of bene�cial elements such as Pt and Y. The advanced coating systems are improved with time and have multifunctional properties such as oxidation resistance, hot corrosion resistance, and thermal barrier.

Because diffusion is a time- and temperature-dependent process, superalloys are not only sub-jected to oxidation but also diffusion of carbon (carburization), nitrogen (nitridation), and so on, depending on service environment. “Mixed attack” (Pettit 1983) refers to the situation whereby two or more reactants attack the alloy simultaneously. Generally, both C and N diffusions cause a




deleterious effect on the mechanical and chemical properties of the alloys. The surface appearance and microstructure of oxidized superalloys are given in Figures 1.48 and 1.49, respectively.

1.5.2 HOT CORROSION

Superalloys are subjected to deposit induced deterioration at presence of impurities, which may come from different sources such as fuel, service atmosphere, and so on at certain range of tem-perature. These impurities deposit on the clean metal surfaces and react with the base metal protec-tive oxide scale and form unwanted low-melting point �ux. This corrosion type propagates faster than oxidation and is called “hot corrosion.” Na, Cl, S, and V are the main elements responsible for hot corrosion. Na and Cl usually come from marine atmosphere (sea salt), V may come from impurities in the fuel and S from both industrial atmosphere and fuel. Gas turbine superalloy parts, which are exposed to the hot gas stream, suffer hot corrosion damages. The main deposit in marine atmosphere is Na2SO4 formed during combustion of the fuel in presence of marine contaminants. Hot corrosion occurs in a temperature range of 650–1000°C. Above these temperatures, salt depos-its volatilize and thus cannot deposit on the metal surface; below these temperatures, the required

FIGURE 1.48 Surface appearance of oxidized Rene 80 alloy (6×).

FIGURE 1.49 Microstructure of the same component, both intergranular oxidation (IGO) and uniform oxidation (gray areas) can be seen in this �gure (200×).



reactions for hot corrosion cannot take place. Although a number of models have been developed, an important reaction series in hot corrosion is given below (Kircher 1989).

Na2SO4 = Na2O + SO3

The reaction products can be both acidic (SO3) and basic (Na2O) and react with the superalloys’ protective oxide scale:

Al2O3 + 3SO3 = 2Al3+ + 3SO42−

Al2O3 + Na2O = 2NaAlO2

There are two stages of hot corrosion: incubation and propagation. In incubation or initiation stage, although the impurities deposit on the metal surface, the behavior is not much different from oxidation. In the propagation stage, however, the protective oxide scale is signi�cantly modi�ed and the alloy beneath the scale is attacked.

Similar to many other chemical reactions, temperature affects hot corrosion rate of superalloys to a great extent. There are two distinctive peaks of the hot corrosion rate as temperature increases, type I and type II. High temperature or type I hot corrosion develops between 800°C and 1000°C and low temperature or type II hot corrosion takes place between 650°C and 850°C. Hot corrosion may occur in superalloys in different ways such as acidic �uxing, basic �uxing or sulfur-induced (sul�dation) (Sims et al. 1987).

Utilization of coatings, selection of proper alloys of high resistance to hot corrosion, use of clean fuels, and �ltering of air wherever possible are the main strategic actions taken to tackle hot corro-sion. Regular engine washing is also very useful to battle hot corrosion in gas turbines. This practice ef�ciently cleans away surface deposits such as salt, sul�des, and so on. Cr is the most bene�cial element in superalloys to increase resistance to hot corrosion. Usually, high-strength nickel-based alloys have lower hot corrosion resistance than low strength ones due to lower chromium content.

1.5.3 FATIGUE

Fatigue is one of the most important damage mechanisms, which should be considered in design stage of superalloy parts such as gas turbine blades and vanes. This mechanism needs cyclic stresses and contains three stages: initiation, propagation, and rupture. In initiation stage, cyclic stress causes crack formation by various mechanisms such as dislocation movement. Cracks usu-ally nucleate on the surfaces of the parts. Stress concentration areas such as tool marks, sharp corners, corrosion pits, scratches, nicks and dents are preferred regions for crack nucleation. Once the crack nucleates, the second stage of the fatigue starts: each load cycle causes a cycle of crack opening, closing, and propagation. A special stress �eld usually a small plastic region surrounded by elastic area is created by cyclic loads at the crack tip. The stress intensity (ΔK) at the crack tip can be calculated to predict propagation rate. Porosities, inclusions, and creep cavities accelerate the propagation. Both stages I and II are more sensitive to defects in precipitation-hardened super-alloys than solid solution alloys. Final rupture takes place as the structure is no longer able to bear the operating loads at a critical fatigue crack length (which decreases the remaining cross section to a critically low area).

Cyclic stress can be mechanical, thermal or a combination of both. Stress versus number of cycles (S–N) and crack propagation rate versus stress intensity factor (da/dN–ΔK) data are usually used to calculate fatigue life. The typical appearance of the typical S–N curve and da/dN–ΔK curve for superalloys can be seen in Figures 1.50 and 1.51, respectively.

Two types of fatigue regime can be de�ned in S–N curves; low cycle fatigue (LCF) and high cycle fatigue (HCF). LCF can be characterized by stresses that are above the elastic limit




(e.g., engine start and stop process of the gas turbine) of the material and fatigue failure life (Nf) is below 105 cycles. HCF usually occurs at stress levels below the elastic limit (e.g., vibrations) and fatigue life is above 107 cycles. Resonant vibrations are the primary source of HCF failure. Da/dN–ΔK curves consist of three regions, second stage of the curve, which is �t power law are the stable crack growth regime.

As superalloys almost always work under corrosive or high-temperature service conditions, fatigue becomes more complex. Oxidation, creep, and corrosion signi�cantly affect fatigue life of component. The fracture propagation path may change into creep cavities where both fatigue and creep conditions are satis�ed. Sometimes bene�cial interrelations may also occur, for example,

0

200

400

600

800

1000

1200

1400

1.00E+03 1.00E+04 1.00E+05 1.00E+06 1.00E+07

Stre

ss (M

Pa)

Number of cycles to fracture—Nf

FIGURE 1.50 Typical S−N curve for superalloys.

1.00E–08

1.00E–07

1.00E–06

1.00E–05

1.00E–04

1.00E–03

1.00E–02

1.00E–01

1.00E+00

1 10 100

Crac

k pr

opog

atio

n am

ount

per

cycl

eda

/dN

(inc

h/cy

cle)

Stress intensity factor range (ΔK–ksi√in)

Region I

Region III

Region II

FIGURE 1.51 Typical da/dN−ΔK curves for superalloys.



oxidation causes crack tip blunting that may effectively retard fatigue crack propagation. The so-called corrosion fatigue occurs when the part is under cyclic stress in a corrosive medium. In this case fatigue crack propagation creates a fresh surface at the crack tip, which is in turn quickly cor-roded by the corrosive medium, thus the fatigue crack propagates faster and more easily. Moreover, corrosion products may cause additional “wedge effect” that generates extra stress in the crack tip region to further speed up crack propagation.

A superalloy’s fatigue behavior depends on various factors such as temperature, grain size, microstructure, service environment, heat treatment condition, frequency, dwell time, amplitude of the cyclic stresses, and so on (Branco 1996). Generally, LCF is more decisive than HCF in compo-nent life. Figure 1.52 shows a fatigue ruptured surface of Hastelloy X superalloy. The curved paral-lel lines or the “striations” reveal the fatigue crack propagation steps.

Some important preventative measures are addressed below to avoid or at least minimize fatigue failure:

• Good knowledge of crack initiation and growth• Careful calculation of stress distribution of the components at the design stage• Minimize stress raisers such as deep tool marks, sharp corners, scratches, nicks, dents and

so on during design, production, usage or maintenance steps• Application of regular nondestructive inspections (NDI) to detect short fatigue cracks

before they reach critical size• Utilization of proper coating in corrosive or high-temperature service conditions• Selection of proper superalloy type that has high resistance to fatigue failure• Complete or partial removal of stress raisers such as scratches, nicks or dents by grinding,

blending, stripping methods during maintenance• Creation of compressive stress on the surface of the part by various methods like shot

peening, laser shot peening to retard crack initiation stage by compensating tensile stress• Removal of corrosion products from the surface regularly by grinding, grit blasting, or

chemical cleaning

1.5.4 CREEP

Creep is one of the important and frequently observed damage mechanisms in superalloys. Three conditions are needed for creep to occur; time, temperature, and stress. Generally, the required temperature to creep should approximately be higher than half of the melting point of

FIGURE 1.52 Scanning electron microscope view of fatigue ruptured surface of nickel-based solid solution Hastelloy X superalloy (250×).




the alloy. The typical creep curve for superalloys under constant load and temperature is given in Figure 1.53.

This curve is plotted as strain against time under constant load and temperature. The curve con-sists of three distinguishable zones. In the �rst zone, creep specimen lengthens elastically just after the application of load. In this zone, the strain rate subsides. The second zone is the stable creep propagation zone, or the steady state creep zone where strain rate remains a minimum constant. In this zone, creep changes the microstructure dramatically by diffusion, microcavities occur in preferred areas such as grain boundaries. Figure 1.54 shows a Zone II creep microstructure of a tur-bine blade of directionally solidi�ed cast nickel-based superalloy. In the tertiary or �nal creep zone, microcavities link and microcracking and �ssuring takes place, giving rise to quick lengthening of specimen and deformation acceleration, which �nally leads to rupture.

Another important curve used to characterize creep life of the components is stress rupture curve. This curve is plotted as stress versus temperature under constant time. The stress rupture curves of various superalloys are given in Figure 1.13.

Cre

ep st

rain

(%)

Time (h)

Region II(stable creep zone)

Region III(unstable creep-rupture zone)

Region I(primary creep zone)

Fracture

FIGURE 1.53 Typical creep curve (time vs. strain) for superalloy.

FIGURE 1.54 Stage II creep specimen microstructure of nickel-based directionally solidi�ed cast superal-loy (arrows indicate grain boundary microcavities) (magni�cation 200×).



1.6 RECENT ADVANCES IN SUPERALLOYS

It seems highly probable that superalloys shall preserve the position of being the most useful materi-als for heat resistant applications in the near future. The main improvement from the service temper-ature point of view in the near past has been the development of SX superalloys. Because these alloys do not need grain boundary strengthener elements, which cause drop of incipient melting tempera-ture, they have higher incipient melting temperature. Thanks to the possibility of higher heat treat-ment temperatures, SX superalloys also have better balanced γ ′ and other phase distribution, which implies better creep, stress rupture resistance, and thermo mechanical fatigue strength. Melting point is the most important limitation as superalloys have already come close to their margin. Superalloys can be used at temperatures very close to their incipient melting point owing to intense research activities in alloy design, process improvements, and so on. Melting point is the genetic property of any alloy system and unavoidably restricts the maximum service temperature of the components. Therefore, future designs with higher working temperature ranges need new materials. The most promising candidates are high tech ceramics such as Si3N4, intermetallic compounds, and composites (Tallan 1989). Due to the production dif�culties, low reproducibility, low reliability, and low fracture toughness properties, these materials need more time to replace superalloys.

Research on superalloys continues. Instead of increasing service temperature, the aim of the research activities is to increase the tensile strength, fatigue strength, creep, stress rupture, and corrosion resistance. Effort also continues to increase reliability of the alloys and reduce cost with process improvements and more sophisticated inspection methods.

The application and development on coatings still continue and greatly increase the capabilities of superalloys. Advanced TBCs protect the superalloys from corrosive service environments and reduce the metal temperatures up to about 150°C (Nicholls et al. 1997). Conventional TBCs usually consist of two layers: MCrAlY (M stands for Ni and/or Co) bond coat plus Y2O3-stabilized ZrO2 top coat and are obtained via air plasma spray (APS) technique. Today, usage of nanoparticles, modi-�cation of chemical composition of the coatings with various rare earth elements, modi�cation of the structure of the coatings with different techniques such as laser glazing, vacuum heat treatment, and so on, utilization of more sophisticated coating techniques such as EB-PVD, vacuum plasma spray, and so on, improved performance of the TBCs (Ma 2006, Morks 2010, Wen 2010, Diltemiz 2011). Developments of high corrosion resistance coatings also continue, amorphous silica is one such example (Donachie and Donachie 2002). More detailed information about coatings applied to the superalloys will be given in another chapter of the book.

REFERENCES

Askeland, D.R. 1985. The Science and Engineering of Materials. PWS Engineering: USA.Branco, C.M., Bryne, J. 1996. Elevated temperature fatigue of In718: Effects of stress ratio and frequency.

AGARD-CP-569 Thermal Mechanical Fatigue of Aircraft Engine Materials Conference Proceedings. NATO: Canada.

Brick, R.M., Pense, A.W., Gordon, R.B. 1977. Structure and Properties of Engineering Materials. McGraw-Hill: USA.

Brooks, C.R. 1982. Heat Treatment, Structure and Properties of Nonferrous Alloys. American Society for Metals: Ohio, USA.

Bryer, T.G., Semiatin, S.L., Vollmer, D.C. 1985. Forging Handbook. American Society For Metals: USA.Campbell, F.C. 2006. Manufacturing Technology for Aerospace Structural Materials. Elsevier: Boston, USA.Davis, F.R. 1997. ASM Specialty Handbook Heat Resistant Materials. ASM International: USA.DeAntonio, D.A. 1991. ASM Handbook Vol. 4 Heat Treating: Heat Treating of Superalloys. ASM International:

USA.Decker, R.F. 1979. Source Book on Materials for Elevated Temperature Applications: Strengthening Mechanisms

in Nickel-Base Superalloys, ed. Bradley, E.F. pp. 275–296. American Society for Metals: Ohio, USA.Diltemiz, S.F. 2010. Thermal and mechanical properties optimization of thermal barrier coatings. PhD disserta-

tion, ESOGU University: Turkey.




Diltemiz, S.F., Kushan, M.C. 2011. The effect of laser glazing on high temperature oxidation resistance of thermal barrier coatings. International Conference on Manufacturing and Industrial Engineering, Turkey.

Donachie, M.J., Donachie, S.J. 2002. Superalloys A Technical Guide. ASM International: USA.Durand-Charre, M. 1997. The Microstructure of Superalloys. CRC press: Amsterdam, Holland.Gell, M., Duhl, D.N. 1985. Advanced High Temperature Alloys, Processing and Properties: The Development

of Single Crystal Superalloy Turbine Blades, eds. Allen S.M., Pelloux R.M., Widmer R. American Society For Metals: USA.

Harris, K., Erickson, G.L., Schwer, R.E. 1990. ASM Handbook Vol. 1. Properties and Selection Irons Steels and High Performance Alloys: Directionally Solidi�ed and Single-Crystal Superalloys. ASM International: USA.

Heubner, U. 1998. Nickel Alloys. Marcel Dekker: USA.Jackman, L.A. 1984. Superalloys Source Book: Forming and Fabrication of Superalloys, ed. Donachie, M.J.

American Society for Metals: USA.Johnston, C.E. 1993. ASM Handbook Vol. 16 Machining: Waterjet/Abrasive Waterjet Machining. ASM

International: USA.Kircher, T.A. 1989. Oxidation, sul�dation and hot corrosion: mechanisms and interrelationships.

AGARD-CP-461 High Temperature Surface Interactions Conference Proceedings. NATO: Canada.Kulkarnı, K.M. 1983, Production to Near Net Shape Source Book: Isothermal Forging-from Research to a

Promising New Manufacturing Technology, eds. Van Tyne C.J, Avitzur B. American Society for Metals: USA.

Ma, W., Gong, S., Xu, H., Cao, X. 2006. The thermal cycling behavior of lanthanum–cerium oxide thermal barrier coating prepared by EB–PVD. Surface & Coatings Technology. 200: 5113–5118.

Mevrel, R., Veyes, J.M. 1989. Effect of protective coatings on mechanical properties of superalloys. AGARD-CP-461 High Temperature Surface Interactions Conference Proceedings. NATO: Canada.

Morks, M.F., Berndt, C.C., Durandet, Y., Brandt, M., Wang, J. 2010. Microscopic observation of laser glazed yttria-stabilized zirconia coatings. Applied Surface Science. 256: 6213–6218.

Nash, P. 1992. ASM Handbook Vol. 3 Alloy Phase Diagrams, p. 692, ed. Baker H. ASM International: USA.Nicholls, J.R., Lawson, K.J., Rickerby, D.S., Morrell, P. 1997. Advanced processing of TBC’s for reduced

thermal conductivity. AGARD-R-823 Thermal Barrier Coating Panel Proceedings. NATO: Denmark.Owczarski, W.A. 1984. Superalloys Source Book: Process and Metallurgical Factors in Joining Superalloys

and Other High Service Temperature Materials, ed. Donachie, M.J. American Society for Metals: USA.Pettit, F.S., Goward, G.W. 1983. Coatings for High Temperature Applications: Oxidation-Corrosion-Errosion

Mechanisms of Environmental Degradation of High Temperature Materials, ed. Lang E. Elsevier: London, England.

Portinha, A., Teixeira, V., Carneiro, J. et al. 2005, Characterization of thermal barrier coatings with a gradient in porosity. Surface & Coatings Technology. 195: 245–251.

Reed, R.C. 2006. The Superalloys Fundamentals and Applications. Cambridge.Sims, C.T., Stoloff, N.S., Hagel, W.C. 1987. Superalloys II. John Wiley & Sons: USA.Smith, W.F. 1993. Structure and Properties of Engineering Alloys. McGraw-Hill: USA.Stefenescu, D.M. 1998. ASM Handbook Vol. 15. Casting: Investment casting. ASM International: USA.Stoloff, N.S. 1990. ASM Handbook Vol. 1. Properties and Selection Irons Steels and High Performance Alloys:

Wrought and P/M Superalloys. ASM International: USA.Tallan, N.M. 1989. Technical evaluation report. AGARD-CP-449 Application of Advanced Material for

Turbomachinery and Rocket Propulsion Conference Proceedings. NATO: UK.Thompson, R.G. 1993. ASM Handbook Vol. 6 Welding Brazing and Soldering: Welding Metallurgy of

Nonferrous High Temperature Materials. ASM International: USA.Tien, J.K., Caul�eld, T. 1989. Superalloys, Supercomposites Superceramics. Academic Press: USA.VerSnyder, F.L. 1982. Gas turbine materials COST 50: Superalloy technology. European Collaborative

Program. Conference Proceedings. Liege.Wen, Ma.W., Dong, H., Guo, H., Gong, S., Zheng, X. 2010. Thermal cycling behavior of La2Ce2O7/8YSZ

double-ceramic-layer thermal barrier coatings prepared by atmospheric plasma spraying. Surface & Coatings Technology. 204: 3366–3370.

Wilde, R.F., Kaufman, M. 1975. High Temperature Metallurgy. Lecture Notes. Small Aircraft Engine Dept: USA.

References

1 Chapter 1 - Superalloys for Super Jobs

Askeland, D.R. 1985. The Science and Engineering ofMaterials. PWS Engineering: USA.

Branco, C.M., Bryne, J. 1996. Elevated temperature fatigueof In718: Effects of stress ratio and frequency.AGARD-CP-569 Thermal Mechanical Fatigue of Aircraft EngineMaterials Conference Proceedings. NATO: Canada.

Brick, R.M., Pense, A.W., Gordon, R.B. 1977. Structure andProperties of Engineering Materials. McGrawHill: USA.

Brooks, C.R. 1982. Heat Treatment, Structure and Propertiesof Nonferrous Alloys. American Society for Metals: Ohio,USA.

Bryer, T.G., Semiatin, S.L., Vollmer, D.C. 1985. ForgingHandbook. American Society For Metals: USA.

Campbell, F.C. 2006. Manufacturing Technology for AerospaceStructural Materials. Elsevier: Boston, USA.

Davis, F.R. 1997. ASM Specialty Handbook Heat ResistantMaterials. ASM International: USA.

DeAntonio, D.A. 1991. ASM Handbook Vol. 4 Heat Treating:Heat Treating of Superalloys. ASM International: USA.

Decker, R.F. 1979. Source Book on Materials for ElevatedTemperature Applications: Strengthening Mechanisms inNickel-Base Superalloys, ed. Bradley, E.F. pp. 275–296.American Society for Metals: Ohio, USA.

Diltemiz, S.F. 2010. Thermal and mechanical propertiesoptimization of thermal barrier coatings. PhD dissertation,ESOGU University: Turkey. Fo c taloging pur ose only

Diltemiz, S.F., Kushan, M.C. 2011. The effect of laserglazing on high temperature oxidation resistance ofthermal barrier coatings. International Conference onManufacturing and Industrial Engineering, Turkey.

Donachie, M.J., Donachie, S.J. 2002. Superalloys ATechnical Guide. ASM International: USA.

Durand-Charre, M. 1997. The Microstructure of Superalloys.CRC press: Amsterdam, Holland.

Gell, M., Duhl, D.N. 1985. Advanced High TemperatureAlloys, Processing and Properties: The Development ofSingle Crystal Superalloy Turbine Blades, eds. Allen S.M.,Pelloux R.M., Widmer R. American Society For Metals: USA.

Harris, K., Erickson, G.L., Schwer, R.E. 1990. ASM HandbookVol. 1. Properties and Selection Irons Steels and HighPerformance Alloys: Directionally Solidi�ed andSingle-Crystal Superalloys. ASM International: USA.

Heubner, U. 1998. Nickel Alloys. Marcel Dekker: USA.

Jackman, L.A. 1984. Superalloys Source Book: Forming andFabrication of Superalloys, ed. Donachie, M.J. AmericanSociety for Metals: USA.

Johnston, C.E. 1993. ASM Handbook Vol. 16 Machining:Waterjet/Abrasive Waterjet Machining. ASM International:USA.

Kircher, T.A. 1989. Oxidation, sul�dation and hotcorrosion: mechanisms and interrelationships. AGARD-CP-461High Temperature Surface Interactions ConferenceProceedings. NATO: Canada.

Kulkarnı, K.M. 1983, Production to Near Net Shape SourceBook: Isothermal Forging-from Research to a Promising NewManufacturing Technology, eds. Van Tyne C.J, Avitzur B.American Society for Metals: USA.

Ma, W., Gong, S., Xu, H., Cao, X. 2006. The thermal cyclingbehavior of lanthanum–cerium oxide thermal barrier coatingprepared by EB–PVD. Surface & Coatings Technology. 200:5113–5118.

Mevrel, R., Veyes, J.M. 1989. Effect of protective coatingson mechanical properties of superalloys. AGARD-CP-461 HighTemperature Surface Interactions Conference Proceedings.NATO: Canada.

Morks, M.F., Berndt, C.C., Durandet, Y., Brandt, M., Wang,J. 2010. Microscopic observation of laser glazedyttria-stabilized zirconia coatings. Applied SurfaceScience. 256: 6213–6218.

Nash, P. 1992. ASM Handbook Vol. 3 Alloy Phase Diagrams, p.692, ed. Baker H. ASM International: USA.

Nicholls, J.R., Lawson, K.J., Rickerby, D.S., Morrell, P.

1997. Advanced processing of TBC’s for reduced thermalconductivity. AGARD-R-823 Thermal Barrier Coating PanelProceedings. NATO: Denmark.

Owczarski, W.A. 1984. Superalloys Source Book: Process andMetallurgical Factors in Joining Superalloys and OtherHigh Service Temperature Materials, ed. Donachie, M.J.American Society for Metals: USA.

Pettit, F.S., Goward, G.W. 1983. Coatings for HighTemperature Applications: Oxidation-Corrosion-ErrosionMechanisms of Environmental Degradation of High TemperatureMaterials, ed. Lang E. Elsevier: London, England.

Portinha, A., Teixeira, V., Carneiro, J. et al. 2005,Characterization of thermal barrier coatings with agradient in porosity. Surface & Coatings Technology. 195:245–251.

Reed, R.C. 2006. The Superalloys Fundamentals andApplications. Cambridge.

Sims, C.T., Stoloff, N.S., Hagel, W.C. 1987. SuperalloysII. John Wiley & Sons: USA.

Smith, W.F. 1993. Structure and Properties of EngineeringAlloys. McGraw-Hill: USA.

Stefenescu, D.M. 1998. ASM Handbook Vol. 15. Casting:Investment casting. ASM International: USA.

Stoloff, N.S. 1990. ASM Handbook Vol. 1. Properties andSelection Irons Steels and High Performance Alloys:Wrought and P/M Superalloys. ASM International: USA.

Tallan, N.M. 1989. Technical evaluation report.AGARD-CP-449 Application of Advanced Material forTurbomachinery and Rocket Propulsion ConferenceProceedings. NATO: UK.

Thompson, R.G. 1993. ASM Handbook Vol. 6 Welding Brazingand Soldering: Welding Metallurgy of Nonferrous HighTemperature Materials. ASM International: USA.

Tien, J.K., Caul�eld, T. 1989. Superalloys, SupercompositesSuperceramics. Academic Press: USA.

VerSnyder, F.L. 1982. Gas turbine materials COST 50:Superalloy technology. European Collaborative Program.Conference Proceedings. Liege.

Wen, Ma.W., Dong, H., Guo, H., Gong, S., Zheng, X. 2010.Thermal cycling behavior of La2Ce2O7/8YSZdouble-ceramic-layer thermal barrier coatings prepared byatmospheric plasma spraying. Surface & CoatingsTechnology. 204: 3366–3370.

Wilde, R.F., Kaufman, M. 1975. High Temperature Metallurgy.Lecture Notes. Small Aircraft Engine Dept: USA.

2 Chapter 2 - Tool Condition Monitoringin Machining Superalloys

1. C.T. Sims, N.S. Stolof, and W.C. Hagel, SuperalloysII—High-Temperature. Materials for Aerospace andIndustrial Power, Wiley, New York, 1987.

2. ASM International Handbook Committee, ASM Handbook,Volume 01—Properties and Selection: Irons, Steels, andHigh-Performance Alloys, ASM International, 1990.

3. D.R. Askeland and P.P. Phulé, The Science andEngineering of Materials, 4th Edition, ThomsonEngineering,Australia, 2002.

4. G.S. Brady, H.R. Clauser, and J.A. Vaccari, MaterialsHandbook, 15th Edition, McGraw Hill DR-52, New York, 2002.

5. M.J. Donachie and S.J. Donachie, Superalloys: ATechnical Guide, 2nd Edition, ASM International, MaterialsPark, OH, 2002.

6. R.C. Reed, The Superalloys—Fundamentals andApplications, Cambridge University Press, New York, 2006.

7. Superalloys: A Primer and History.

8. H.K.D.H. Bhadeshia, Nickel Based Superalloys.http://www.msm.cam.ac.uk/phase-trans/2003/Superalloys/superalloys.html (accessed on May 24, 2011).

9. M. McLean, Nickel-base superalloys: Current status andpotential, in: High-Temperature Structural Materials, ed.RW Cahn, AG Evans and M. McLean. Chapman and Hall,London,1996.

10. ASM Specialty Handbook: Heat Resistant Materials, ASMInternational, May 1, 1997.

11. ASM Specialty Handbook: Nickel, Cobalt, and TheirAlloys, ASM International, 2000.

12. R.E. Smallman and R.J. Bishop, Modern PhysicalMetallurgy and Materials Engineering, 6th Edition,Butterworth-Heinemann, London, 1999. For catalogingpurposes only

13. E.O. Ezugwu, J. Bonney, and Y. Yamane, An overview ofthe machinability of aeroengine alloys, Journal ofMaterial Processing Technology, 134, 233–253, 2003.

14. E.O. Ezugwu, Key improvements in the machining ofdif�cult-to-cut aerospace superalloys, InternationalJournal of Machine Tools & Manufacture, 45, 1353–1367,2005.

15. T.H.C. Childs, K. Maekawa, T. Obikawa, and Y. Yamane,Metal Machining—Theory and Applications, Arnold, London,2000.

16. G.T. Smith, Cutting Tool Technology: IndustrialHandbook, Springer, London, 2008.

17. Machining Nickel Alloys. A Nickel Development InstituteReference Book, Series No 11 008. http://www.

18. E.M. Trent and P.K. Wright, Metal Cutting, 4th Edition.Butterworth–Heinemann, Boston, 2000.

19. G. Schneider Jr., Cutting Tool Applications, GMRSAssociates, 2002.

20. R. Arunachalam and M.A. Mannan, Machinability ofnickel-based high temperature alloys, Machining Scienceand Technology, 4(1), 127–168, 2000.

21. H. Tschätsch, Applied Machining Technology, Springer,New York, 2009.

22. ASM Handbook Volume 16: Machining, ASM International,1989.

23. F.C. Campbell, Manufacturing Technology for AerospaceStructural Materials, Elsevier Ltd., 2006.

24. Machining-Special Metals Corporation Products.http://www.specialmetals.com/documents/machining.pdf(accessed on June 2011)

25. M. Ay, U. Ulas¸, and A. Hasçalik, Effect of traversespeed on abrasive waterjet machining of age hardenedInconel 718 nickel-based superalloy, Materials andManufacturing Processes, 25(10), 1160–1165, 2010.

26. H.Z. Li, X.Q. Chen, H. Zeng, and X.P. Li, Embedded toolcondition monitoring for intelligent machining,International Journal of Computer Applications inTechnology, 28(1), 74–81, 2007.

27. A.M. Fong, X.Q. Chen, and H.Z. Li, Overview of material

processing automation, in: X.Q. Chen, R. Devanathan, andA.M. Feng (Eds.), Advanced Automation Techniques inAdaptive Material Processing, World Scienti�c PublishingCo. Pte. Ltd., New Jersey, 1–18, 2002.

28. R. Komanduri and T.A. Schroeder, On shear instabilityin machining a Nickel-Iron base superalloy, ASME Journalof Engineering for Industry, 108, 93–100, 1986.

29. S. Hanasaki, J. Fujiwara, and M. Touge, Tool wear ofcoated tools when machining a high nickel alloy, Annals ofthe CIRP, 39/1, 77–80, 1990.

30. G. Byrne and B. Bienia, Tool life scatter when millingwith TiN-coated HSS indexible inserts, Annals of the CIRP,40/1, 45–48, 1991.

31. E.O. Ezugwu, K.A. Olajire, and A. Jawaid, Wearperformance of multiplayer-coated carbide tools, MachiningScience and Technology, 5/1, 115–119, 2001.

32. M. Alauddin, M.A. Mazid, M.A. El Baradi, and M.J.S.Hashmi, Cutting forces in the end milling of Inconel 718,Journal of Materials Processing Technology, 77, 153–159,1998.

33. E.-G. Ng, D.W. Lee, A.R.C. Sharman, R.C. Dewes, andD.K. Aspinwall, High speed ball nose end milling of Inconel718, Annals of the CIRP, 49/1, 41–46, 2000.

34. H.Z. Li, H. Zeng, and X.Q. Chen, An experimental studyof tool wear and cutting force variation in the endmilling of Inconel 718 with coated carbide inserts, Journalof Materials Processing Technology, 180(1–3), 296–304,2006.

35. H.Z. Li, A. Albrecht, and X.Q. Chen, A tool wearobserver model for �ank wear monitoring in the milling ofnickel-based alloys, Special Issue on “Smart MachiningSystems,” International Journal of Mechatronics andManufacturing Systems (IJMMS), 2(5/6), 620–637, 2009.

36. X.Q. Chen and H.Z. Li, Development of a tool wearobserver model for online tool condition monitoring andcontrol in machining nickel-based alloys, InternationalJournal of Advanced Manufacturing Technology, 45(7–8),786–800, 2009.

37. H. Weber and T.N. Loladze, Grundlagen des Spanens, VEBVerlag Technik, Berlin, 1986.

38. J.-J. Park and A.G. Ulsoy, On-line tool wear estimationusing an adaptive observer and computer vision, Part 1:Theory, Journal of Engineering for Industry, Transactionsof the ASME, 115, 30–36, 1993.

39. J.-J. Park and A.G. Ulsoy, On-line tool wear estimationusing an adaptive observer and computer vision, Part 2:Experiment, Journal of Engineering for Industry,Transactions of the ASME, 115, 37–43, 1993.

40. J.-J. Park and A.G. Ulsoy, On-line tool wear estimationusing force measurement and a non-linear observer, Journalof Engineering for Industry, Transactions of the ASME, 114,666–672, 1992.

41. Y. Koren, A.G. Ulsoy, and K. Danai, Tool wear andbreakage detection using a process model, Annals of theCIRP, 35/1, 283–288, 1986.

42. Y. Koren, Flank wear model of cutting tools usingcontrol theory, Journal of Engineering for Industry,Transactions of the ASME, 100, 103–109, 2000.

43. A. Novak and H. Wiklund, On-line prediction of the toollife, Annals of the CIRP, 45/1, 93–96, 1996.

44. K. Danai and A.G. Ulsoy, An adaptive observer foron-line tool wear estimation in turning, Part I: Theory,Mechanical Systems and Signal Processing, l(2), 211–225,1987.

45. K. Danai and A.G. ULSOY, An adaptive observer foron-line tool wear estimation in turning, Part II: Results,Mechanical Systems and Signal Processing, l(2), 227–240,1987.

46. E. Kuljanic and M. Sortino, TWEM, a method based oncutting forces—Monitoring tool wear in face milling,International Journal of Machine Tools & Manufacture, 45,29–34, 2005.

47. D. Shi and N.N. Gindy, Tool wear predictive model basedon least squares support vector machines, MechanicalSystems and Signal Processing, 21(4), 1799–1814, 2007.

48. A. Iqbal, N. He, N.U. Dar, and L. Li, Comparison offuzzy expert system based strategies of of�ine and onlineestimation of �ank wear in hard milling process, ExpertSystems with Applications, 33, 61–66, 2007. For cataloging

purposes only

3 Chapter 3 - Laser Cladding and Alloyingfor Aerospace Applications

1. Gnanamuthu, D. S. 1976. Cladding, US, Patent 3942180.

2. Gnanamuthu, D. S. 1980. Laser surface treatment. Opt.Eng. 19:783–792.

3. Weerasinghe, V. M. and Steen, W. M. 1983. Laser claddingby powder injection. In Transport Phenomena in MaterialsProcessing, J. M. M. M. Chen and C. Tucker, Eds. (ASME, NewYork), pp. 15–23.

4. Lin, J. and Steen, W. M. 1998. An in-process method forthe inverse estimation of the powder catchment ef�ciencyduring laser cladding. Opt. Laser Technol. 30:77–84.

5. Li, L., Steen, W. M., and Hibbero, R. D. 1990.In-process clad quality monitoring using optical method. InProceedings of the Conference on Laser-Assisted Processing,Boston, MA, USA, Vol. 1279, pp. 89–100.

6. Weerasinghe, V. M. and Steen, W. M. 1983. Computersimulation model for laser cladding. Am. Soc. Mech. Engrs.10:15–23.

7. Ellis, M., Xiao, D. C., Lee, C., Steen, W. M., Watkins,K. G., and Brown, W. P. 1995. Processing aspects of lasercladding an aluminium alloy onto steel. J. Mater. Process.Technol. 52:55–67.

8. Jeng, J. Y., Quayle, B. E., Modern, P. J., Steen, W. M.,and Bastow, B. D. 1993. Laser surface treatments toimprove the intergranular corrosion resistance of 18/13/Nband 304L in nitric acid. Corros. Sci. 35:1289–1293.

9. Kar, A. and Mazumder, J. 1988. One-dimensional�nite-medium diffusion model for extended solid solution inlaser cladding of Hf on nickel. Acta Metal. 36:701–702.

10. Zhong, M. and Liu, W. 2010. Laser surface cladding, thestate of the art and challenges. Proceedings of iSME SpeIssue 224:1041–1060.

11. Agrawal, G., Kar, A., and Mazumder, J. 1993.Theoretical studies on extended solid solubility andnonequilibrium phase diagram for Nb-Al alloy formed duringlaser cladding. Scr. Metall. Mater. 28:1453–1458.

12. Sircar, S., Singh, J., and Mazumder, J. 1989.

Microstructure evolution and nonequilibrium phase diagramfor Ni-Hf binary alloy produced by laser cladding. ActaMetal. 37:1167–1176.

13. Mazumder, J. and Kar, A. 1987. Solid solubility inlaser cladding. J. Met. 39:18–23.

14. Kar, A. and Mazumder, J. 1987. Effect of cooling rateon solid solubility in laser cladding. ASME 3:237–249.

15. Mazumder, J. and Singh, J. 1986. Laser surface alloyingand cladding for corrosion and wear. NATO ASI E. Appl.Sci. 7:297–307.

16. Tuominen, J., Vuoristo, P., Manttyla, T., Latokartano,J., Vihinein, J., and Andersson, J. 2003. Microstructureand corrosion behavior of high power diode laser depositedInconel 625 coating. J. Laser Appl. 15:55–61.

17. Ribaudo, C. and Mazumder, J. 1989. Oxidation behaviorof a laser-clad nickel-based alloy containing hafnium.Mater. Sci. Eng. A 121–122:531–538.

18. Singh, J. and Mazumder, J. 1986. In-situ formation ofNi-Cr-Al-R. E. alloy by laser cladding with mixed powderfeed. In Proceedings of the Third International Conferenceon Lasers in Manufacturing, SpringerVerlag, pp. 169–179.

19. Dinda, G. P., Dasgupta, A. K., and Mazumder, J. 2009.Laser aided direct metal deposition of Inconel 625superalloy: Microstructural evolution and thermalstability. Mater. Sci. Eng. A 509:98–104.

20. Mazumder, J., Dutta, D., Kikuchi, N., and Ghosh, A.2000. Closed loop direct metal deposition: Art to part.Opt. Lasers Eng. 34:397–414.

21. Vilar, R. 1999. Laser cladding. J. Laser Appl. 11:64–79.

22. Li, W. B., Engstrom, H., and Powell, J. 1995. Modelingof the laser cladding process-preheating of the blownpowderc material. Lasers Eng. 4:329–341.

22. Salehi, D. and Brandt, M. 2006. Melt pool temperaturecontrol using LabVIEW in Nd:YAG laser blown powdercladding process. Int. J. Adv. Manuf. Technol. 29:273–278.

23. Syed, W. U. H., Pinkerton, A., Liu, Z., and Li, L.2007. Coincident wire and powder deposition by laser toform compositionally graded material. Surf. Coat. Technol.

201:7083–7091.

24. Pinkerton, A., Syed, W. U. H., and Li, L. 2007.Theoretical analysis of the coincident wire-powder laserdeposition process. ASME Trans. J. Manuf. Sci. Eng.129:1019–1027.

25. Powell, J., Henry, P. S., and Steen, W. M. 1988. Lasercladding with preplaced powder: Analysis of thermalcycling and dilution effects. Surf. Eng. 4:141–149.

26. Huang, S. H., Nolan, D., and Brandt, M. 2003. Preplaced WC/Ni clad layers produced with a pulsed Nd:YAGlaser via optical �bres. Surf. Coat. Technol. 165:26–34.

27. Guo, L. F., Yue, T. M., and Man, H. C. 2004. A �niteelement method approach for thermal analysis of lasercladding of magnesium alloy with preplaced Al-Si powder. J.Laser Appl. 16:229–235.

28. Lin, J. and Steen, W. M. 1998. Design characteristicsand development of a nozzle for coaxial laser cladding. J.Laser Appl. 10:55–63.

29. Pawlowski, L. 1999. Thick laser coatings: A review. J.Thermal Spray Technol. 8:279–295.

30. Dubourg, L. and Archambeault, J. 2008. Technologicaland scienti�c landscape of laser cladding process in 2007.Surf. Coat. Technol. 202:5863–5869.

31. Jeng, J.Y. and Lin, M. C. 2001. Mold fabrication andmodi�cation using hybrid processes of selective lasercladding and milling. J. Mater. Process. Technol.110:98–103.

32. Zhou, S. F., Huang, Y. J., and Zeng, X. Y. 2008. Astudy of Ni basedWC composite coatings by laser inductionhybrid rapid cladding with elliptical spot. Appl. Surf.Sci. 254:3110–3119.

33. Lin, J. M. 2002. Process monitoring and �ow control incoaxial laser cladding. Adv. Laser Opt. Res. 1:103–125.

34. Salehi, D., Brandt, M., and Kogel-Hollacher, M. 2002.Fundamental investigations on process monitoring in thelaser cladding process. In Proceedings of ICALEO 2002,Scottsdale, AZ, USA, 14–17, pp. 599–611.

35. Tucker, T. R., Clauer, A. H., and Wright, I. G. 1984.

Laser processed composite metal cladding for slurryerosion resistance. Thin Solid Films 118:73–84.

36. Vandehaar, E., Molian, P. A., and Baldwin, M. 1988.Laser cladding of thermal barrier coatings. Surf. Eng.4:159–172.

37. Langer, S., Weisheit, A., and Mordikem, B. L. 1998.Laser cladding of an AlSiBi-alloy for bearing applications.Lasers Eng. 8:35–41.

38. Peng, Z. X. and Boyer, R. E. 2001. High-power CO 2laser cladding of WC based particles for cutting toolapplications. In Proceedings of ICALEO 2001, USA, 92–93:pp. 594–602.

39. Marchione, T. 2002. Industrial applications of lasercladding. In Proceedings of ICALEO 2002, Scottsdale, AZ,USA, pp. 631–636. For c t l in urposes o ly

40. Jendrzejewski, R., Sliwinski, G., and Conde, A. 2003.Direct laser cladding of Ni-based alloy powder forindustrial applications. In Proceedings of 14thInternational Symposium on Gas ©ow, Chemical Lasers, andHigh-Powerlasers, Wroclaw, Poland, Vol. 5120: pp. 688–691.

41. Goswami,G. L., Kumar, S., and Galun, R. 2003. Lasercladding of Ni-Mo alloys for hardfacing applications.Lasers Eng. 13:1–12.

42. Jendrzejewski, R., Sliwinski, G., and Conde, A. 2003.Laser cladding of Ni- and Co-based coatings for turbineindustry applications. Laser Technol. VII: Appl. Lasers5229:233–238.

43. Smurov, I. 2008. Laser cladding and laser assisteddirect manufacturing. Surf. Coat. Technol. 202:4496–4502.

44. Nowotny, S., Scharek, S., and Beyer, E. 2007. Laserbeam build-up welding: precision in repair, surfacecladding and direct 3D metal deposition. J. Therm. SprayTechnol. 16:344–348.

45. Dinda, G. P., Dasgupta, A. K., and Mazumder, J. 2009.Laser aided direct metal deposition of Inconel 625superalloy: Microstructural evolution and thermalstability. Mater. Sci. Eng. A 509:98–104.

46. Choi, J. and Hua, Y. 2004. Dimensional and materialcharacteristics of direct deposited H13 tool steel by CO 2

laser. J. Laser Appl. 16:245–251.

47. Zhong, M. L., Liu, W. J., and Ning, G. Q. 2004. Laserdirect manufacturing of tungsten nickel collimationcomponent. J. Mater. Process. Technol. 147:167–173.

48. Mazumder, J., Schifferer, A., and Choi, J. 1999. Directmaterials deposition: designed macro and microstructure.Mater. Res. Innov. 3:118–131.

49. Pinkerton, A. J. and Li, L. 2003. The effect of laserpulse width on multiple-layer 316L steel cladmicrostructure and surface �nish. Appl. Surf. Sci.208–209:411–416.

50. Pinkerton, A. J. and Li, L. 2004. The behaviour ofwater and gas-atomised tool steel powders in coaxial laserfreeform fabrication. Thin Solid Films 453–454:600–605.

51. Pinkerton, A. J. and Li, L. 2005. An experimental andnumerical study of the in�uence of diode laser beam shapeon thin wall direct metal deposition. J. Laser Appl.17:47–56.

52. Pinkerton, A. J. and Li, L. 2005. Direct additive lasermanufacturing using gas- and water-atomised H13 tool steelpowders. Int. J. Adv. Manuf. Technol. 25:471–479.

53. Syed,W. U. H. and Li, L. 2005. Effects of wire feedingdirection and location in multiple layer diode laserdirect metal deposition. Appl. Surf. Sci. 248:518–524.

54. Syed, W. U. H., Pinkerton, A., and Li, L. 2006.Combining wire and coaxial powder feeding in laser directmetal deposition for rapid prototyping. Appl. Surf. Sci.252:312–317.

55. Steen, W. M. 1998. Laser Material Processing, Springer.(With kind copyright permission from Springer Science+Business Media).

56. Pawlowski, L. 1999. Thick laser coatings: A review. J.Thermal Spray Technol. 8(2):279–295.

57. Gedda, H., Kaplan, A., and Powell, J. 2005. Melt–solidinteractions in laser cladding and laser casting. Metall.Mater. Trans. B 36B:683–689.

58. Folkes, J. A. 1994. Developments in laser surfacemodi�cation and coating. Surf. Coat. Technol. 63:65–71.

59. Possin, G.E., Parks, H.G., and Chiang, S.W. 1981.Convection in Pulsed Laser Formed Melts, Laser andElectron-Beam Solid Interactions and Materials Processing,J.F. Gibbons, L.D. Hess, and T.W. Sigmon, Eds., Elsevier,North-Holland, pp. 73–80.

60. Almeida, A., Anjo, S. M., Vilar, R., Li, R., Fereira,M. G. S., Steen, W. M., and Watkins, K. G. 1995. Laseralloying of aluminum alloys with chromium. Surf. Coat.Technol. 70:221–229.

61. Carvalho, P.A. and Vilar, R. 1997. Laser alloying ofzinc with aluminum: Solidi�cation structures. Surf. Coat.Technol. 91:158–166.

62. Ignatiev, M., Kovalev, E. I., Melekhin, I., Smurov, Y.,and Sturlese, S. 1993. Investigation of the hardening of atitanium alloy by laser nitriding. Wear 166:233–236.

63. Gerdes, C. Karimi, A., and Bieler, H. W. 1995. Waterdroplet erosion and microstructure of laser-nitridedTi–6Al–4V. Wear 186–187:368–374.

64. Graydon, O. 1998. Jets of molten metal make industrialparts. Opt. Lasers Eng. 33–35.

65. Kurz, W., Bezencon, C., and Gaumann, M. 2001. Columnarto equiaxed transition in solidi�cation processing. Sci.Technol. Adv. Mater. 2:185–191.

66. Schneider, M. 1998. Thesis on “Laser Cladding withPowder” University of Twente, Netherland, open literature.

67. Picasso, M. and Hoadley, A. F. A. 1994. Finite elementsimulation of laser surface treatments includingconvection in the melt pool. Int. J. Num. Meth. Heat FluidFlow. 4:61–83.

68. Wu, X., Zhu, B., Zeng, X., Hu, X., and Cui, K. 1996.Critical state of laser cladding with powder autofeeding.Surf. Coat. Technol. 79:200–204.

69. Gu, Yu., Changsheng, D., Minlin, Z., Mingxing, M., Lin,L., and Wenjin, L. 2011. Fabrication of nanoporousmanganese by laser cladding and selective electrochemicalde-alloying. Appl. Surf. Sci. 257:3211–3215.

70. Santhanakrishnan, S., Kong, F., and Kovacevic, R. 2011.An experimentally based thermo-kinetic hardening model for

high power direct diode laser cladding. J. Mater. Process.Technol. 211:1247–1259.

71. Majumdar, J. D. and Manna, I. 2011. Laser materialprocessing. Int. Mater. Rev. 56(5/6):341–388.

72. Frenk, A. and Kurz, W. 1993. High speed laser cladding:Solidi�cation conditions and microstructure of acobaltbased alloy. Mater. Sci. Eng. A 173, 339–342.

73. Mei, Z., Guo, L. F., and Yue, T. M. 2005. The effect oflaser cladding on the corrosion resistance of magnesiumZK60/SiC composite. J. Mater. Proces. Technol. 161:462–466.

74. Osório, W.R., Cheung, N., Spinelli, J.E., Cruz, K.S.,and Garcia, A. 2008. Microstructural modi�cation by lasersurface remelting and its effect on the corrosionresistance of an Al–9 wt%Si casting alloy. Appl. Surf.Sci. 254:2763–2770.

75. Pariona, M. M., Teleginski, V., Santos, K. D., Machado,S., Zara, A. J., Zurba, N. K., and Riva, R. 2012. Yb-�berlaser beam effects on the surface modi�cation of Al–Feaerospace alloy obtaining weld �let structures, low �neporosity and corrosion resistance. Surf. Coatings Technol.206(8–9):2293–2301.

76. Green, M. and Sharp, M. 2010. Report by Association ofLaser Users (AILU) for the Photonics and PlasticElectronics Knowledge Transfer Network.

77. Kahlen, F. J., Von Klitzing, A., and Kar, A. InProceedings of the ICALEO 99, Vol. 87, Laser Institute ofAmerica, San Diego, California, November 1999, pp.F129–F137.

78. Grif�th, M.L., Harwell, L.D., Romero, J.A., Schlienger,E., Atwood, C.L., and Smugeresky, J.E. 1997. Proceedingsof the Solid Freeform Fabrication Symposium, Austin, TX, p.387.

79. Brooks, J. A., Robino, C. V., Grif�th, M. L., Headley,T. J., Yang, N. C. Y., Goods, S. H., Susan, D. F., andPuskar, J. D., LDRD 2001. Final Report: Exploiting LENSTechnology Through Novel Materials. Sandia ReportSAND2001–8186.

80. Lin, X. and Yue, T. M. 2005. Phase formation andmicrostructure evolution in laser rapid forming of gradedSS316L/Rene88DT alloy. Mater. Sci. Eng. A 402:295–306.

81. Overtone, G. 2009. Materials processing: Laser additivemanufacturing gains strength. Laser Focus World, 6.

82. Fortuna, D. 2002. Gas turbine repairs with binder freebraze alloy. Sulzer Tech. Rev. 4:18–20.

83. Martein H. and Van Dijk, H. 2006. Repairing Aero-engineparts, Industrial Laser Solutions, 21(5).

84. Cao, X., Jahazi, M., Immarigeon, J. P., and Wallace, W.2006. A review of laser welding techniques for magnesiumalloys. J. Mater. Process. Technol. 171:188–204.

85. Barrallier, L., Fabre, A., Masse, J. E., and Ceretti,M. 2002. Residual stress measurements using neutrondiffraction in magnesium alloy laser welded joints. Mater.Sci. Forum 404–407:399–404.

86. Lehner, C., Reinhart, G., and Schaller, L. 1999.Welding of die cast magnesium alloys for production. J.Laser Appl. 11(5):206–210.

87. Klausecker, R., Geiger, M., Otto, A., and Haberling, C.2001. Laser beam joining of magnesium matrix compositeswith magnesium alloys. In Proceedings of the LANE 2001(3rd): Laser Assisted Net Shape Engineering, pp. 323–332,Vol. 3.

88. Haferkamp, H., Bach, F. W., Burmester, I., Kreutzburg,K., and Niemeyer, M. 1996. Nd:YAG laser beam welding ofmagnesium constructions. In Proceedings of the ThirdInternational Magnesium Conference, UMIST, Manchester, UK,pp. 89–98.

98. Sanders, P. G., Keske, J. S., Leong, K. H., andKornecki, G. 1999. High power Nd:YAG and CO 2 laserwelding of magnesium. J. Laser Appl. 11(2):96–103.

90. Leong, K. H., Kornecki, G., Sanders, P. G., and Keske,J. S. 1998. Laser beam welding of AZ31B-H24 alloy, ICALEO98: Laser Materials Processing Conference, Orlando, FL,16–19 November, pp. 28–36.

91. Marya, M. and Edwards, G. R. 2000. The laser welding ofmagnesium alloy AZ91. Weld. World 44(2):31–37.

92. Akman, E., Demir, A., Canel, T., and Sınmazçelik, T.2009. Laser welding of Ti6Al4V titanium alloys.J. Material. Process. Technol. 209:3705–3713.

93. Kreimeyer, M., Wagner, F., and Vollertsen, F. 2005.Laser processing of aluminum–titanium-tailored blanks.Opt. Lasers Eng. 43:1021–1035.

94. Bechert, D. W., Bruse, M., and Hage, W. 2000.Experiments with three-dimensional riblets as an idealizedmodel of shark skin. Exp. Fluids 28(5):403–412.

95. Woodcoc, D. A., Lightfoot, P., and Ritter, C. 2000.Negative thermal expansion in Y 2 (WO 4 ) 3 . J. SolidState Chem. 149:92–98.

96. Sumithra, S. and Umarji, A. M. 2005. Hygroscopicity andbulk thermal expansion in Y 2 W 3 O 12 . Mater. Res. Bull.40(1):167–176.

97. Schaaf, P., Schikor, H., Höche, D., Lange, C.,Drescher, V., and Wilden, J. 2008. Laser clad surfaces forshark-skin effect by high-temperature activation. Surf.Coating Technol. 203:470–475. F r cataloging pu p e only

98. Wynholds, H. 1994. A Survey of Waterproo�ng Proceduresfor Shuttle. NASA Report, Contract A39157DC(BCV). Reportprepared for the Thermal Protection Branch, NASA, AmesResearch Center, Moffett, Field, CA.

99. Goldstein, H. 1982. Fibrous Ceramic Insulation (AmesResearch Center). Presented at NASA/AIAA AdvancedMaterials Echnology Conference, No. 16–17.

100. Zorba, V., Tzanetakis, P., Fotakis, C., Spanakis, E.,Stratakis, E., Papazoglou, D. G., and Zergioti, I. 2006.Silicon electron emitters fabricated by ultraviolet laserpulses. Appl. Phys. Lett. 88:081103–081105.

101. Her, T. H., Finlay, J., Wu, C., Deliwala, S., andMazur, E. 1998. Microstructuring of silicon withfemtosecond laser pulses. Appl. Phys. Lett. 73:1673–1675.

102. Zorba, V., Persano, L., Pisignano, D., Athanassiou,A., Stratakis, E., Cingolani, R., Tzanetakis, P., andFotakis, C. 2006. Making silicon hydrophobic: Wettabilitycontrol by two-length scale simultaneous patterning withfemto second laser irradiation. Nanotechnology17:3234–3238.

103. Bauerle, D. 2000. Laser Processing and Chemistry,Springer, Berlin, Heidelberg.

104. Murthy, N. S., Prabhu, R. D., Martin, J. J., Zhou, L.,and Headrick, R. L. J. 2006. Self-assembled and etchedcones on laser ablated polymer surfaces. Appl. Phys. A100:023528–023528-12.

105. Bensaoula, A., Boney, C., Pillai, A., Shafeev, G. A.,Simakin, A. V., and Starikov, D. 2004. Etch characteristicsof GaN and BN materials in chlorine-based plasmas. Appl.Phys. A 79:973–975.

106. Barberoglou, V. Z., Stratakis, E., Spanakis, E.,Tzanetakis, P., Anastasiadis, H., and Fotakis, C. 2009.Bio-inspired water repellent surfaces produced by ultrafastLaser structuring of silicon. Appl. Surf. Sci.255:5425–5429.

107. Morvan, P. 2000. Filters for the aerospace industry.News from Prospace 40–41.

108. Madhavan, P. 1988. Advances in aircraft transmissionlubricant �ltration technology. AIAA 1988–2984.

109. Trego, L. 1997. Hydraulic system puri�cation.Aerospace Eng. 17:31–34.

110. Jen, C. 2000. A high capacity advanced fuel �ltrationsystem. ADA384193, MSI-TR-8002.

111. Multi-phase method for evaluating �ltrationperformance of �ne lube �lter elements utilized inaerospace power and propulsion lubrication systems. SAEStandard ARP5454, 2003.

112. Pajak, P. T. 1991. Laser machining: Theory andPractice. Springer-Verlag, New York.

113. Garnero, P., Jamin, A., and Lecardonne, L. 2004.Bi-propellant propulsion system improvement fortelecommunication satellites. AIAA-2004-4211.

114. Weia, W., Dia, Z., Allen, D. M., and Almond, H. J. A.2008. Non-traditional machining techniques for fabricatingmetal aerospace �lters. Chin. J. Aeronaut. 21:441–447.

115. David, S. A., Vitek, J. M., Babu, S. S., Boatner, L.A., and Reed, R. W. 1997. Welding of nickel basesuperalloy single crystals. Sci. Technol. Weld Joining2:79–88.

116. Vitek, J. M., David, S. A., and Boatner, L. A. 1997.

Microstructural development in single crystal nickel basesuperalloy welds. Sci. Technol Weld Joining 2:109–118.

117. Liu, W. and Dupont, J. N. 2004. Effects of melt-poolgeometry on crystal growth and microstructure developmentin laser surface-melted superalloy single crystals.Mathematical modeling of single-crystal growth ina meltpool (part I). Acta Material. 52:4833–4847.

118. Bansal, R., Acharya, R., Gambone, J. J., and Das, S.2011, Experimental and theoretical analysis of scanninglaser epitaxy applied to nickel based superalloys. AugProceedings International Symposium, Solid PreformFabrication, University of Texas at Austin.

119. Wilfried, K. 2011. Laser repair and cladding ofsuperalloy single crystals. Int. Symposium on Joining andSustain of Superalloys M &Ms T, October 16–20, 2011,Columbus, OH.

4 Chapter 4 - High-PerformanceWear/Corrosion-Resistant Superalloys

1. J. R. Davis, Cobalt-base alloys, in Nickel, Cobalt, andTheir Alloys, ASM International, Materials Park, 2000.

2. A. Frenk and J. D. Wagnière, Laser cladding withcobalt-based hardfacing, J. Phys. IV France, 1(C7), 1991,65–68.

3. W. Betteridge, Cobalt and Its Alloys, Halsted Press,Chichester, 1982.

4. A. Aizaz and P. Kumar, Properties of Stellite alloy No.21 made via pliable powder technology, Met. Powd. Rep.,40(9), 1985, 507–510.

5. A. Davin and D. Coutsouradis, Development of abrasion-and corrosion-resistant alloys for use in aqueous media,Cobalt, 52, 1971, 160–161.

6. C. T. Sims, A contemporary view of cobalt-base alloys,JOM, 21, 1969, 27–42.

7. K. Ando, T. Omori, J. Sato, Y. Sutou, K. Oikawa, R.Kainuma and K. Ishida, Effect of alloying elements onfcc/hcp martensitic transformation and shape memoryproperties in Co-Al alloys, Mater. Trans., 47(9), 2006,2381–2386.

8. I. Campos, G. Ramirez, U. Figueroa and C. Velazquez,Paste bonding process: Evaluation of boron mobility onborided steels, Surf. Eng., 23(3), 2007, 216-222.

9. R. Liu and D. Y. Li, Protective effect of yttriumadditive in lubricants on corrosive wear, Wear, 225-229,1999, 968–974.

10. D. Raghu and J. B. Wu, Recent development in wear andcorrosion-resistant alloys, Mater. Perform., 36(11), 1997,27–36.

11. A. Halstead and R. D. Rawlings, Effect of ironadditions on the microstructure and properties of the“Tribaloy” Co-Mo-Cr-Si wear resistant alloys, J. Mater.Sci., 20(5), 1985, 1693–1704.

12. Anon, Unique intermetallics combat wear and corrosion,Metal. Prog., 107(5), 1975, 90-92.

13. C. B. Cameron, R. A. Hoffman and R. W. Poskitt,Tribaloy intermetallic alloy compositions—New materials oradditives for wear resistant applications, Prog. Powd.Metall., 31, 1975, 41–51.

14. S. E. Mason and R. D. Rawlings, Structure and hardnessof Ni-Mo-Cr-Si wear and corrosion resistant alloys, Mater.Sci. Technol., 5(2), 1989, 180–185.

15. S. C. Lim and M. F. Ashby, Wear-mechanism maps, ActaMetall., 35(1), 1987, 1–24.

16. A. Frenk and W. Kurz, Microstructural effects on thesliding wear resistance of a cobalt-based alloy, Wear,174(1-2), 1994, 81–91.

17. C. J. Heathcock, A. Ball and B. E. Protheroe,Cavitation erosion of cobalt-based Stellite alloys,cemented carbides and surface-treated low alloy steels,Wear, 74(1), 1981, 11–26.

18. W. L. Silence, Effect of structure on wear resistanceof Co-, Fe-, and Ni-base, in K. C. Ludema and S. K. Rhee(Ed.), Proceedings of the International Conference on Wearof Materials, ASME, New York, 1987, pp. 77–85.

19. L. C. Wang and D. Y. Li, Effects of yttrium onmicrostructure, mechanical properties, high-temperaturewear behaviour of cast Stellite 6 alloy, Wear, 255(1–6),2003, 535–544.

20. I. Radu, D. Y. Li and R. Llewellyn, Tribologocalbehaviour of Stellite 21 modi�ed with yttrium, Wear,257(11), 2004, 1154–1166.

21. I. Radu and D. Y. Li, The wear performance ofyttrium-modi�ed Stellite 712 at elevated temperatures,Tribol. Int., 40(2), 2007, 254–265.

22. H. Celik and M. Kaplan, Effects of silicon on the wearbehaviour of cobalt-based alloys at elevated temperature,Wear, 257(5–6), 2004, 606–611.

23. R. D. Schmidt and D. P. Ferriss, New materialsresistant to wear and corrosion to 1000°C, Wear, 32(3),1975, 279–289.

24. W. L. Wilson, T. A. Galioto, R. L. Miller and G. A.Whitlow, Tribological and corrosion behavior of materialsfor high temperature sodium service, Micro. Sci., 7, 1979,

265–274.

25. R. B. Herchenroeder, S. J. Matthew, J. W. Tackett andS. T. Wlodek, HAYNES alloy No. 188, Cobalt, 54, 1972,3–13.

26. D. L. Klarstrom, Wrought cobalt-base superalloys, J.Mater. Eng. Perform., 2(4), 1993, 523–530.

27. G. Y. Lai, J. J. Barnes and J. E. Barnes, A burner riginvestigation of the hot corrosion behaviour of severalwrought superalloys and intermetallics, ASME, 91-GT-21,1991, 1–10.

28. B. A. Boeck, T. H. Sanders, Jr., V. Anand, A. J. Hickland P. Kumar, Relationships between processing,microstructure, and tensile properties of a Co-Cr-Mo alloy,Powd. Metall., 28(2), 1985, 97–104.

29. G. A. Whitlow, J. C. Cwynar, R. L. Miller and S. L.Schrock, Sodium corrosion behavior of alloys for fastreactor applications, in Chemical Aspects of Corrosion andMass Transfer in Liquid Sodium, Metall. Soci. AIME, NewYork, 1973, 1–63.

30. Y. D. Zhang, Z. G. Yang, C. Zhang and H. Lan, Oxidationbehavior of Tribaloy T-800 alloy at 800 and 1000°C, Oxid.Metal., 70, 2008, 229–239.

31. J. L. De Mol Van Otterloo and J. T. M. DeHosson,Microstructural features and mechanical properties of acobalt-based laser coating, Acta Mater., 45(3), 1997,1225–1236.

32. Y. Birol, Thermal fatigue testing of Inconel 617 andStellite 6 alloys as potential tooling materials forthixoforming of steels, Mater. Sci. Eng. A, 527(7–8), 2010,1938–1945.

33. C. B. Cameron and D. P. Ferriss, Tribaloy intermetallicmaterials: New wear- and corrosion-resistant alloys,Anti-Corr. Meth. Mater., 22(4), 1975, 5–8.

34. G. K. Bansal, Effect of �aw shape on strength ofceramics, J. Am. Ceram. Soc., 59(1-2), 1976, 87–88.

35. W. L. Karlsen, Effect of nitrogen alloying on Stellite21 cobalt chromium molybdenum biomedical implant alloy:Processing and microstructure, Mater. Sci. Forum, 318-320,1999, 303–308.

36. D. H. E. Persson, S. Jacobson and S. Hogmark, Thein�uence of phase transformations and oxidation on thegalling resistance and low friction behavior of a laserprocessed Co-based alloy, Wear, 254(11), 2003, 1134–1140.

37. P. Huang, R. Liu, X. J. Wu and M. X. Yao, Effects ofmolybdenum content and heat treatment on mechanical andtribological properties of a low-carbon Stellite alloy, J.Eng. Mater. Technol., 129(4), 2007, 523–529.

38. Y. Ning, P. C. Patnaik, R. Liu, M. X. Yao and X. J. Wu,Effects of fabrication process and coating ofreinforcements on the microstructure and wear performanceof Stellite alloy composites, Mater. Sci. Eng. A,391(1-2), 2005, 313–324.

39. C. D. Opris, R. Liu, M. X. Yao and X. J. Wu,Development of Stellite alloy composites with sintering/HIPing technique for wear-resistant applications, Mater.Desi., 28, 2007, 581–591.

40. S. A. Tsukerman, Powder Metallurgy, Pergamon Press,Oxford, 1965.

41. R. Yosomiya, K. Morimoto, A. Nakajima, Y. Ikada and T.Suzuki, Adhesion and Bonding in Composites, Marcel DekkerInc., New York, 1990.

42. S. H. Kang, T. Shinoda, T. Kato and H. S. Jeong,Thermal fatigue characteristics of PTA hardfaced steels,Surf. Eng., 17(6), 2001, 498–504.

43. R. Liu, M. X. Yao, P. C. Patnaik and X. J. Wu, Animproved wear-resistant PTA hardfacing—VWC/ Stellite 21, J.Comp. Mater., 40(24), 2006, 2203–2215.

44. W. H. Jiang, X. D. Yao, H. R. Guan and Z. Q. Hu,Carbidebehaviourduringhightemperaturelowcyclefatigueinacobalt-basesuperalloy, J. Mater. Sci., 34(12), 1999,2859–2864.

45. W. H. Jiang, X. D. Yao and Z. Q. Hu, Secondary M 6 Cprecipitation in a cobalt-base superalloy, J. Mater. Sci.Lett., 18(4), 1999, 303–305.

46. Zoltek Committee, User’s Guide for Short Carbon FiberComposites, Zoltek Companies Inc., USA, 2000.

47. W. Xu, R. Liu, P. C. Patnaik, M. X. Yao and X. J. Wu,

Mechanical and tribological properties of newly developedTribaloy alloys, Mater. Sci. Eng. A, 452–453, 2007,427–436.

48. R. Liu, W. Xu, M. X. Yao, P. C. Patnaik and X. J. Wu, Anewly developed Tribaloy alloy with increased ductility,Scripta Material., 53(12), 2005, 1351–1355.

49. R. Liu, M. X. Yao, P. C. Patnaik and X. J. Wu,Investigation of mechanical behavior of cobalt-basedintermetallic materials using a nano-indentation technique,J. Advan. Mater., Special Edition 1(2), 2007, 65–75.

50. A. Halstead and R. D. Rawlings, The fracture behaviorof two Co-Mo-Cr-Si wear resistant alloys (“Tribaloy”), J.Mater. Sci., 20(4), 1985, 1248–1256.

51. M. X. Yao, J. B. C. Wu, W. Xu and R. Liu,Microstructural characteristics and corrosion resistance inmolten Zn-Al bath of Co-Mo-Cr-Si alloys, Mater. Sci. Eng.A, 407(1–2), 2005, 299–305.

52. M. X. Yao, J. B. C. Wu, S. Yick, Y. S. Xie and R. Liu,High temperature wear and corrosion resistance of a newlydeveloped Laves phase strengthened Co-Mo-Cr-Si alloy,Mater. Sci. Eng. A, 435–436, 2006, 78–83.

53. H. Inaba and H. Yukawa, Correlation between theparameters obtained from the DV-Xα molecular orbitalcalculation and the partial molar enthalpies of formationfor liquid Ni-,Fe-,Cr- and Ti-based transition metalalloys, J. Alloys Comp., 248, 1997, 168–175. For catalogingpu po es only

54. A. Halstead and R. D. Rawlings, Structure and hardnessof Co-Mo-Cr-Si wear resistant alloys (Tribaloys), MetalSci., 18(10), 1984, 491–500.

55. R. Liu, X. S. Qi, S. Kapooir and X. J. Wu,Investigation of solidi�cation behaviour and associatemicrostructures of Co-Cr-W and Co-Cr-Mo alloy systems usingDSC technique, J. Mater. Sci., 45(22), 2010, 6225–6234.

56. R. Liu, X. S. Qi, S. Kapooir and X. J. Wu, Effects ofchemical composition on solidi�cation, microstructure andhardness of Co-Cr-W-Ni and Co-Cr-Mo-Ni alloy systems, Int.J. Res. Rev. Appl. Sci., 5(2), 2010, 110–122.

57. H. Baker, Introduction to phase diagrams, alloy phasediagrams, in ASM Handbook, Vol. 3, ASM International,

Materials Park, 1992.

58. J. R. Davis, Metals Handbook, ASM International,Materials Park, 1998.

59. Y. N. Liu, H. Yang, G. Tan, S. Miyazaki, B. Jiang andY. Liu, Stress-induced FCC ↔ HCP martensitictransformation in CoNi, J. Alloy. Comp., 368(1-2), 2004,157–163.

60. W. F. Brown, Jr. and J. E. Strawley, Plane Strain CrackToughness Testing of High Strength Metallic Materials,ASTM STP, No. 410, American Society for Testing andmaterials, USA, 1966.

61. D. M. Buckley, Adhesion, friction and wear of cobaltand cobalt-base alloys, Cobalt, 38, 1968, 20–28.

62. R. Kamel and K. Halim, The effect of phase change onthe mechanical properties of cobalt near its transformationtemperature, Phys. Status Solidi., 15, 1966, 63–69.

63. V. Kuzucu, M. Ceylan, H. Celik and I. Aksoy,Microstructure and phase analyses of Stellite 6 plus 6 wt.%Mo alloy, J. Mater. Proc. Technol., 69, 1997, 257–263.

64. K. Rajan, Thermodynamics assessment of heat treatmentsfor a Co-Cr-Mo alloy, J. Mater. Sci., 18, 1983, 257–264.

65. M. X. Yao, J. B. C. Wu, W. Xu and R. Liu,Metallographic study and wear resistance of a high-Cwrought Co-based alloy Stellite 706 K, Mater. Sci. Eng. A,407(1–2), 2005, 291–298.

66. W. C. Oliver and G. M. Pharr, An improved technique fordetermining hardness and elastic modulus using load anddisplacement sensing indentation experiments, J. Mater.Res., 7, 1992, 1564–1583.

5 Chapter 5 - High-Temperature Oxidationof Aerospace Materials

1. Ohnabe H, Wasaki S, Imamura R, Status and outlook ofhigh temperature materials applied to aeroengines, inProceedings of the Third International Symposium onUltra-high Temperature Materials, Tajimi, 1993, pp.112–126. Based on D. W. Petraset et al., Metal Progress,1986.

2. Antolovich S D, Stusrud R W, MacKay R A, Anton D L, KhanT, Kissinger R D, Klarstrom D L, Eds. Superalloys. TheMineral, Metals and Materials Society, Warrendale, PA,1992.

3. Zhao J C, Westbrook J H, Ultra high temperaturematerials for jet engines. MRS Bulleting 28, 2003,622–627.

4. Dimiduk D M, Miracle D B, Ward C H, Development ofintermetallic materials for aerospace systems. MaterialsScience and Technology 8, 1992, 367–375.

5. Dimiduk D M, Gamma titianium aluminides—An emergingmaterials technology, in Gamma Titanium Aluminides, Y. W.Kim, R. Wagner, M. Yamaguchi, Eds. The Mineral, Metals andMaterials Society, Warrendale, PA, p. 3, 1995.

6. Yamaguchi M, Inui H, TiAl Compounds for structrualapplication, in Structural Intermetallics, R. Darolia,J.J. Lewandowski, C. T. Liu, P. L. Martin, D. B. Miracle,M. V. Nathal, Eds. The Mineral, Metals and MaterialsSociety, Warrendale, PA, p. 127–142, 1993.

7. Wagner R, Appel F, Dogan B, Ennis P J, Lorenz U,Mtillauer L, Nicolai H P et al. Investment casting ofgamma-TiAl-based alloys: Microstructure and data base forgas turbine applications, in Gamma Titanium Aluminides, Y.W. Kim, R. Wagner, M. Yamaguchi, Eds. The Mineral, Metalsand Materials Society, Warrendale, PA, p. 387–404, 1995.

8. Liu C T, George E P, Sikka V K, Deevi S C, Design of Ni3 Al alloys for structural use, in Processing and DesignIssues in High Temperature Materials, N. S. Stoloff, R. H.Jones, Eds. The Mineral, Metals and Materials Society,Warrendale, PA, pp. 139–155, 1996.

9. Morsi K, The diversity of combustion synthesisprocessing: A review. Materials Science and Engineering A299, 2001, 1–15.

10. Chan K S, Theoretical analysis of grain size effects ontensile ductility. Scripta Metallurgica et Materialia 24,1990, 1725–1730.

11. Ward-Close C M, Minor R, Doorbar P J,Intermetallic-matrix composites—A review. Intermetallics 4,1996, 217–229.

12. Noebe R D, Walston W S, Prospects for development ofstructural NiAl alloys, in Structural Intermetallics, M.V. Nathal, R. Darolia, C. T. Liu, P. L. Martin, D. B.Miricale, R. Wagner, M. Yamaguchi, Eds. The Minerals,Metals and Materials Society, Warrendale, PA, pp. 573–584,1997.

13. Appel F, Wagner R, Microstructure and deformation oftwo-phase γ-titanium aluminides. Materials Science andEngineering: R: Reports 22, 1998, 187–268.

14. Larsen J M, Russ S M, Jones J W, An evaluation of�ber-reinforced titanium matrix composites for advancedhigh-temperature aerospace applications. Metallurgical andMaterials Transactions A— Physical Metallurgy and MaterialsScience 26, 1995, 3211–3223.

15. Dressler W, Riedel R, Progress in silicon-basednon-oxide structural ceramics. International Journal ofRefractory Metals and Hard Materials 15, 1997, 13–47.

16. Evans A G, Blumenthal W I, High temperature failure inceramics, in Fracture Mechanics of Ceramics (Vol. 5), R.C. Bradt, A. G. Evans, D. P. H. Hasselman, F. F. Lange,Eds. Plenum Press, New York, pp. 423–448, 1983.

17. Smith J T, Quackenbusch C L, Phase effects in Si 3 N 4containing Y 2 O 3 or CeO 2 . American Ceramic SocietyBulletin 59, 1980, 529–532.

18. Giachello A, Martinengo P C, Tommasini G, Sintering andproperties of silicon-nitride containing Y 2 O 3 and MgO.American Ceramic Society Bulletin 59, 1980, 1212–1215.

19. Wiebke G, Die oxidation von siliziumkarbid. Berichteder Deatschen Keramischen Gesellschaft 37, 1960, 219–226.

20. Schlichting J, Schwetz K, Oxidation behaviour of densesintered alpha-silicon carbide. High Temperatures HighPressures 14, 1982, 219–223.

21. Grathwohl G, Reets T, Thummler F, Creep of hot-pressedand sintered SiC with different sintering additives.Science of Ceramics 11, 1981, 425–431.

22. Yao Z, Stiglich J, Sudarshan T S, Molybdenum silicidebased materials and their properties. Journal of MaterialsEngineering and Performance 8, 1999, 291–304.

23. Mitra R, Mechanical behavior and oxidation resistanceof structural silicides. International Materials Reviews51, 2006, 13–64.

24. Fitzer E, Maurer H J, Diffusion and precipitationphenomena in aluminized and chromium-aluminized iron- andnickel-base alloys, in Materials and Coaitngs to ResistHigh Temperature Corrosion, D. R. Holmes, A. Rahmel, Eds.Applied Science, London, U.K., pp. 253–269, 1978.

25. Restall J E, Malik M, Singheiser L, Metallic diffusionand overlay coatings, in Proceeding of High TemperatureAlloys for Gas Turbines and Other Applications, W. Betz,Ed. D. Reidel Publishing Co., Dordrecht, 1986, p357.

26. Goward G W, Progress in coatings for gas turbineairfoils. Surface Coating and Technology 108–109, 1998,73–79.

27. Nicholls J R, Simms N J, Chan W Y, Evans H E, Smartoverlay coatings—Concept and practice. Surface andCoatings Technology 149, 2002, 236–244.

28. Nicholls J R, Hancock P, Al-Yasiri L H, Optimisingoxidation resistance of MCrAl coating systems using vapourphase alloy design. Materials Science and Technology 5,1989, 799–805.

29. Peng X, Wang F, Oxidation resistant nanocrystallinecoatings, in New Developments in High TemperatureCorrosion, W. Gao, Z. Li, Eds. Woodhead Publishing Limited,Cambridge, U.K., pp. 456–475, 2008.

30. Wang F, Lou H, Zhu S, Wu W, The mechanism of scaleadhesion on sputtered microcrystallized CoCrAl �lms.Oxidation of Metals 45, 1996, 39–50. Fo catalogingpurposes only

31. Liu Z, Gao W, Dahm K, Wang F, Oxidation behavior ofsputter-deposited Ni-Cr-Al microcrystalline coatings. ActaMaterialia 46, 1998, 1691–1700.

32. Waki H, Kobayashi A, Development of lateral compressionmethod of circular tube thin coating for mechanicalproperties of plasma sprayed CoNiCrAlY. MaterialsTransactions 47, 2006, 1626–1630.

33. Brandl W, Toma D, Kruger J, Grabke H J, Matthaus G, Theoxidation behaviour of HVOF thermalsprayed MCrAlY coatings.Surface and Coatings Technology 94–95, 1997, 21–26.

34. Hirose A, Kotani H, Kobayashi K, Surface modi�cation bylaser cladding of Ni-Cr-Al-Y alloy. Tetsu toHagane—Journal of the Iron and Steel Institute of Japan 79,1993, 105–112.

35. Padture N P, Gell M, Jordan E H, Thermal barriercoatings for gas-turbine engine applications. Science 296,2002, 280–284.

36. Clarke D R, Levi C G, Materials design for the nextgeneration thermal barrier coatings. Annual Review ofMaterials Research 33, 2003, 383–417.

37. Kamaraj M, Rafting in single crystal nickel-basesuperalloys—An overview. Sadhana-Academy Proceedings inEngineering Sciences 28, 2003, 115–128.

38. Koizumi Y, Kobayashi T, Jianxing Z, Yokokawa T, HaradaH, Aoki Y, Arai M, Development of a nextgeneration Ni-basesingle crystal superalloy, in Proceeding of theInternational Gas Turbine Congress, Tokyo, 2003, pp. 1–6.

39. Nazmy M, Epishin A, Link T, Staubli M, A review ofdegradation in single crystal nickel based superalloys.Energy Materials 1, 2006, 263–268.

40. Sato A, Yeh A-C, Kobayashi T, Yokokawa T, Harada H,Murakumo T, Zhnag J X, Fifth generation Ni based singlecrystal superalloy with superior elevated temperatureproperties, Energy Materials 2, 2007, 19–25.

41. Giggins C S, Pettit F S, Oxidation of Ni-Cr-Al alloysbetween 1000 and 1200°C. Journal of the ElectrochemicalSociety 118, 1971, 1782–1790.

42. Kear B H, Pettit F S, Fornwalt D E, Lemaire, On thetransient oxidation of a Ni-15Cr-6Al alloy. Oxidation ofMetals 3, 1971, 557–569.

43. Wagner C, Passivity and inhibition during the oxidationof metals at elevated temperatures. Corrosion Science 5,

1965, 751–764.

44. Wagner C, Reaktionstypen bei der oxydation vonlegierungen. Zeitschrift fur Elektrochemie 63, 1959,772–790.

45. Jacob K T, Electrochemical determination of activitiesin Cr 2 O 3 -Al 2 O 3 solid solution. Journal of theElectrochemical Society 125, 1978, 175–179.

46. Kim S S, Sanders T H, Thermodynamic modeling of theisomorphous phase diagrams in the Al 2 O 3 -Cr 2 O 3 and V2 O 3 -Cr 2 O 3 systems. Journal of American CeramicSociety 84, 2001, 1881–1884.

47. Niu Y, Zhang X J, Wu Y, Gesmundo F, The third-elementeffect in the oxidation of Ni-xCr-7Al(x = 0,5,10,15 at%)alloys in 1atm O 2 at 900–1000°C. Corrosion Sciences 48,2006, 4020–4036.

48. Rapp R A, The transition from internal to externaloxidation and the formation of interruption bands insilver-indium alloys. Acta Metallurgica 10, 1961, 730–741.

49. Bowman R, Superalloys: A Primer and History, asupplement to The Minerals, Metals & Materials Society’ssite dedicated to The 9th International Symposium onSuperalloys.

50. Pfeil L B, Improvement in heat-resisting alloys, UKPatent, 1937, No. 459848.

51. Jacobson N S, Myers D L, Zhu D, Humphrey D L,Rhenium/oxygen interactions at elevated temperatures.Oxidation of Metals 55, 2001, 471–480.

52. Duriez C, Rhenium oxidation by steam at hightemperatures. Oxidation of Metals 61, 2004, 49–67.

53. Moniruzzaman M, Maeda M, Murata Y, Morinaga M,Degradation of high-temperature oxidation resistance forNi-based alloys by Re addition and the optimization ofRe/Al content. ISIJ International 43, 2003, 386–393.

54. Akhtar A, Hegde S, Reed R C, The oxidation ofsingle-crystal nickel-based superalloys. JOM 58, 2006,37–42.

55. Akhtar A, Hook M S, Reed R C, On the oxidation of thethird-generation single-crystal superalloy CMSX-10.

Metallurgical and Materials Transactions A 36A, 2005,3001–3017.

56. Göbel M, Rahmel A, Schütze M, The isothermal-oxidationbehavior of several nickel-base single-crystal superalloyswith and without coatings. Oxidation of Metals 39, 1993,231–261.

57. Aoki Y, Arai M, Research and Development of NewGeneration Single Crystal Superalloy for HPT Blades, inProceedings of the 2nd Symposium of EnvironmentallyCompatible Propulsion System for Next GenerationSupersonic Transport, Tokyo, 2004, paper F-1.

58. Sato A, Harada H, Kobayashi T, Murakumo T, Zhang J X,Yokokawa T, A 5th generation Ni-base single crystalsuperalloy designed for the combination of excellentoxidation resistance and creep strength at elevatedtemperatures. Journal of the Japan Institute of Metals 70,2006, 196–199.

59. Kawagishi K, Harada H, Sato A, Sato A, Kobayashi T, Theoxidation properties of fourth generation single-crystalnickel-based superalloys. JOM, 58, 2006, 43–46.

60. Tomaszewicz P, Wallwork G R, Iron-aluminum alloys: Areview of their oxidation behavior. Review HighTemperature Materials 4, 1978, 75–105.

61. Stott F H, Wood G C, Hobby M G, A comparison of theoxidation behavior of Fe-Cr-Al, Ni-Cr-Al, and Co-Cr-Alalloys. Oxidation of Metals 3, 1971, 103–113.

62. Kofstad P, High Temperature Corrosion, Elsevier AppliedScience, New York, 1988, p. 373.

63. Whittle D P, Stringer J, Improvements in hightemperature oxidation resistance by additions of reactiveelements or oxide dispersions. Philosophical Transactionsof the Royal Society of London Series A— MathematicalPhysical and Engineering Sciences 295, 1980, 309–329.

64. Moon D P, Role of reactive elements in alloyprotection. Materials Science and Technology 5, 1989,754–764.

65. Pieraggi B, Rapp R A, Chromia scale growth in alloyoxidation and the reactive element effect. Journal of theElectrochemistry Society 140, 1993, 2844–2850.

66. Hou P Y, Stringer J, The effect of reactive elementadditions on the selective oxidation, growth and adhesionof chromia scales. Materials Science and Engineering: A202, 1995, 1–10.

67. Stott F H, Wood G C, Stringer J, The in�uence ofalloying elements on the development and maintenance ofprotective scales. Oxidation of Metals 44, 1995, 113–145.

68. Pint B A, Experimental observation in support of thedynamic-segregation theory to explain the reactiveelementeffect. Oxidation of Metals 45, 1996, 1–37.

69. Peng X, Guan Y, Dong Z, Xu C, Wang F, A fundamentalaspect of the growth of alumina scale on a metal withdispersion of CeO 2 nanoparticles. Corrosion Science 53,2011, 1954–1959.

70. Ramanarayanan T A, Ayer R, Petkovic-Luton R, Leta D P,The in�uence of yttrium on oxide scale growth andadherence. Oxidation of Metals 29, 1988, 445–472.

71. Weinbruch S, Anastassiadis A, Ortner H M, Martinz H P,Wilhartitz, On the mechanism of high temperature oxidationof ODS superalloys: Signi�cance of yttrium depletion withinthe oxide scales. Oxidation of Metals 51, 1999, 111–128.

72. Pint B A, Martin J R, Hobbs L W, The oxidationmechanism of theta-Al 2 O 3 scales. Solid State Ionics 78,1995, 99–107.

73. Graham M J, Eldridge J I, Mitchell D F, Hussey R J,Anion transport in growing Cr 2 O 3 and Al 2 O 3 scales.Materials Science Forum 43, 1989, 207–242.

74. Tien J K, Pettit F S, Mechanism of oxide adherence onFe-25Cr-4Al (Y or Sc) alloys. Metallurgical Transactions3, 1972, 1587–1599.

75. Giggins C S, Kear B H, Pettit F S, Tien J K, Factorsaffecting adhesion of oxide scales on alloys.Metallurgical Transactions 5, 1974, 1685–1688.

76. Hindam H, Whittle D P, Mechanism of PEG growth andin�uence on scale adhesion, in Proceeding of JapanInstitute of Metal, 3rd International Symposium on HighTemperature Corrosion of Metals, Lake Yamanaka, Japan,1982, p. 261.

77. Hutchings R, Loretto M H, Smallman R E, Oxidation of

intermetallic compound NiAl. Metal Science 15, 1981, 7–13.

78. Rapp R A, The high-temperature oxidation of metalsforming cation-diffusion scales. MetallurgicalTransactions A—Physical Metallurgy and Materials Science15, 1984, 765–782.

79. Brumm M W, Grabke H J, Oxidation behavior of NiAl-II.Cavity formation beneath the oxide scale on NiAl ofdifferent stoichiometries. Corrosion Science 34, 1993,547–553.

80. Peng X, Wang F, Morphologic investigation and growth ofthe alumina scale on magnetron-sputtered CoCrAl NCs withand without yttrium. Corrosion Science 45, 2003, 2293–2306.

81. Funkenbusch A W, Smeggil J G, Bornstein N S, Reactiveelement—Sulfur interaction and oxide scale adherence.Metallurgical Transactions A—Physical Metallurgy andMaterials Science 16, 1985, 1164–1166.

82. Lees D G, On the reasons for the effects of dispersionsof stable oxides and additions of reactive elements on theadhesion and growth-mechanisms of chromia and aluminascales-the sulfur effect. Oxidation of Metals 27, 1987,75–81.

83. Hou P Y, Segregation phenomena at thermally grown Al 2O 3 /alloy interfaces. Annual Review of Materials Research38, 2008, 275–298.

84. Pettit F S, Oxidation mechanisms for nickel-aluminumalloys at temperatures between 900 and 1300°C.Transactions of the Metallurgical Society of AIME 239,1967, 1296.

85. Schumann E, Rühle M, Microstructural observations onthe oxidation of γ ′—Ni 3 Al at high oxygen partialpressure. Acta Metallurgica et Materialia 42, 1994,1481–1487. For cataloging pur oses ly

86. Stoloff N S, Liu C T, Deevi S C, Emerging applicationsof intermetallics. Intermetallics 8, 2000, 1313–1320.

87. Taniguchi S, Shibata T, Cyclic oxidation behavior of Ni3 Al—0.1B base alloys containing a Ti, Zr, or Hf addition.Oxidation of Metals 25, 1986, 201–216.

88. Doychak J, Smialek J L, Mitchell T E, Transientoxidation of single-crystalβ-NiAl. Metallurgical

Transactions A—Physical Metallurgy and Materials Science20, 1989, 499–518.

89. Schumann E, The effect of Y-ion implantation on theoxidation of beta-NiAl. Oxidation of Metals 43, 1995,157–172.

90. Doychak J, Rühle M, TEM studies of oxidized NiAl and Ni3 Al cross sections. Oxidation of Metals 31, 1989,431–452.

91. Peng X, Li T, Pan W P, Oxidation of a La 2 O 3 -modi�edaluminide coating. Scripta Materialia 44, 2001, 1033–1038.

92. Kuenzly J D, Douglass D L, Oxidation mechanism of Ni 3Al containing yttrium. Oxidation of Metals 8, 1974,139–178.

93. Choi S C, Cho H J, Kim Y J, Lee D B, High temperatureoxidation behavior of pure Ni 3 Al. Oxidation of Metals46, 1996, 51–72.

94. Pint B A, Hobbs L W, The oxidation behavior of Y 2 O 3-dispersed β-NiAl, Oxidation of Metals 61, 2004, 273–292.

95. Peng X, Li M, Wang F, A novel ultra�ne-grained Ni 3 Alwith increased cyclic oxidation resistance. CorrosionScience 53, 2011, 1616–1620.

96. Choudhury N S, Graham H G, Hinze J W, Oxidationbehavior of titanium aluminides, in Properties of HighTemperature Alloys with Emphasis on Environmental Effects,Z. A. Foroulis, F. S. Pettit, Eds. The ElectrochemicalSociety, Inc., Princeton, NJ, p. 668–680, 1976.

97. Perkins R A, Chiang K T, Meier G H, Formation ofalumina on TiAl alloys. Scripta Metallurgica 21, 1987,1505–1510.

98. Meier G H, Fundamentals of the oxidation of hightemperature intermetallics, in Oxidation of HighTemperatureIntermetallics, T. Grobstein, J. Doychak, Eds. TheMinerals, Metals and Materials Society, Warrendale, PA, p.1, 1989.

99. Luthra K L, Stability of protective oxide �lms onTi-base alloys. Oxidation of Metals 36, 1991, 475–490.

100. Becker S, Rahmel A, Schorr M, Schütze M, Mechanism ofisothermal oxidation of the intermetallic TiAl and of TiAl

alloys. Oxidation of Metals 38, 1992, 425–464.

101. Kobayashi E, Yoshihara M, Tanaka R, Improvement inoxidation resistance of the intermetallic compound titaniumaluminide by heat-treatment under a low partial-pressureoxygen atmosphere. High Temperature Technology 8, 1990,179–184.

102. Fergus J F, Review of the effect of alloy compositionon the growth rates of scales formed during oxidation ofgamma titanium aluminide alloys. Materials Science andEngineering A 338, 2002, 108–125.

103. Subrahmanyam J, Cyclic oxidation of aluminizedTi-14Al-24Nb alloy. Journal of Materials Science 23, 1988,1906–1910.

104. Figge U, Elschner A, Zheng N, Schuster H, Quadakkers WJ, Surface analytical investigations on the oxidationbehavior of TiAl-base intermetallics. Fresenius Journal ofAnalytical Chemistry 346, 1993, 75–78.

105. Shida Y, Anada H, The in�uence of ternary elementaddition on the oxidation behaviour of TiAl intermetalliccompound in high temperature air. Corrosion Science 35,1993, 945–953.

106. Maki K, Shioda M, Sayashi M, Shimizu T, Isobe S,Effect of silicon and niobium on oxidation resistance ofTiAl intermetallics. Materials Science and Engineering A153, 1992, 591–596.

107. Shemet V, Tyagi A K, Becker J S, Lersch P, SingheiserL, Quadakkers W J, The formation of protectivealumina-based scales during high-temperature air oxidationof gamma-TiAl alloys. Oxidation of Metals 54, 2000,211–235.

108. Varma S K, Chan A, Mahapatra R N, Static and cyclicoxidation of Ti-44Al and Ti-44Al-xNb alloys. Oxidation ofMetals 55, 2001, 423–435.

109. Kim B G, Kim G M, Kim C J, Oxidation behavior ofTiAl-X(X = Cr, V, Si, Mo or Nb) intermetallics at elevatedtemperature. Scripta Metallurgical et Materialia 33, 1995,1117–1125.

110. Kekare S A, Aswath P B, Oxidation of TiAl basedintermetallics. Journal of Materials Science 32, 1997,2485–2499.

111. Haanappel V A C, Sunderkötter J D, Stroosnijder M F,The isothermal and cyclic high temperature oxidationbehaviour of Ti-48Al-2Mn-2Nb compared with Ti-48Al-2Cr-2Nband Ti-48Al-2Cr. Intermetallics 7, 1999, 529–541.

112. Stroosnijder M F, Zheng N, Quadakkers W J, Hoffman R,Gil A, Lanza F, The effect of niobium ion implantation onthe oxidation behavior of a gamma-TiAl-based intermetallic.Oxidation of Metals 46, 1996, 19–35.

113. Pérez P, Adeva P, Improvement of oxidation behavior ofa Ti-48Al-2Cr alloy by a nitridation treatment. Oxidationof Metals 56, 2001, 271–285.

114. Lu W, Chen C L, He L L, Wang F H, Lin J P, Chen G L,(S)TEM study of different stages ofTi-45Al8Nb-0.2W-0.2B-0.02Y alloy oxidation at 900°C.Corrosion Science 50, 2008, 978–988.

115. Taniguchi S, Uesaki K, Zhu Y-C, Matsumoto Y, ShibataT, In�uence of implantation of Al, Si, Cr or Mo ions onthe oxidation behaviour of TiAl under thermal cycleconditions. Materials Science and Engineering A 266, 1999,267–275.

116. Singheiser L, Quadakkers W J, Shemet V, Isothermal andcyclic oxidation behavior of titanium aluminides, in the2nd International Symposium on Gamma Titanium AluminidesHeld in Conjunction with the TMS 1999 Annual Meeting, SanDiego, CA, 1999, pp. 743–752.

117. Shemet V, Tyagi A K, Penkalla J, Singheiser L, BeckerJ S, Quadakkers W J, Mechanisms of scale formation on newalumina forming gamma-TiAl alloys studied by MCs + -SIMS.Materials at High Temperatures 17, 2000, 41–47.

118. Narita T, Izumi T, Yatagai M, Yoshioka T, Sul�dationprocessing and Cr addition to improve oxidation resistanceof TiAl intermetallics in air at 1173K. Intermetallics 8,2000, 371–379.

119. Schumacher G, Dettenwanger F, Schütze M,IberlA,ReilD,XRD investigations of the phosphorus effecton the oxidation of TiAl alloys by high-resolution methods.Oxidation of Metals 54, 2000, 317–337.

120. Schütze M, Hald M, Improvement of the oxidationresistance of TiAl alloys by using the chlorine effect.Materials Science and Engineering A 239–240, 1997, 847–858.

121. Leyens C, Braun R, Fröhlich M, Hovsepian P E, Recentprogress in the coating protection of gammatitanium-aluminides. JOM 58, 2006, 17–21.

122. Narushima T, Goto T, Hirai T, Iguchi Y,High-temperature oxidation of silicon carbide and siliconnitride. Materials Transactions JIM 38, 1997, 821–835.

123. Aldinger F, Auweter-Kurtz M, Fertig M, Herdrich G,Hirsch K, Linder P, Matusch D, Neuer G, Schumacher U,Winter M, Behaviour of reusable heat shield materials underRE-Entry conditions, in Basic Research and Technologiesfor Two-Stage-to-Orbit Vehicles, D. Jacob, G. Sachs, S.Wagner, Eds. Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim,Germany, pp. 527–548, 2005.

124. Infed F, Weiland S, Lange H, Steinacher A, Handrick K,Presented at 6th European Workshop on Thermal Protectionand Hot Structures, Stuttgart, April, 2009.

125. Jacobson N S, Corrosion of silicon-based ceramics incombustion environments. Journal of the American CeramicSociety 76, 1993, 3–28.

126. Jacobson N S, Myers D L, Active oxidation of SiC.Oxidation of Metals 75, 2011, 1–25.

127. Wagner C, Passivity during the oxidation of silicon atelevated temperatures. Journal of Applied Physics 29,1958, 1295–1297.

128. Turkdogan E T, Grievson P, Darken L S, Enhancement ofdiffusion-limited rates of vaporization of metals. Journalof Physical Chemistry 67, 1963, 1647–1654.

129. Balat M, Flamant G, Male G, Pichelin G, Active topassive transition in the oxidation of silicon-carbide athigh-temperature and low-pressure in molecular and atomicoxygen. Journal of Materials Science 27, 1992, 697–703.

130. Balat M, Determination of the active-to-passivetransition in the oxidation of silicon carbide in standardand microwave-excited air. Journal of the European CeramicSociety 16, 1996, 55–62.

131. Nickel K G, The role of condensed silicon monoxide inthe active-to-passive oxidation transition of siliconcarbide. Journal of the European Ceramic Society 9, 1992,3–8.

132. Lou V L K, Mitchell T E, Heuer A H, Graphical displaysof the thermodynamics of high-temperature gas-solidreactions and their application to oxidation of metals andevaporation of oxides. Journal of the American CeramicSociety 68, 1985, 49–58.

133. Heuer A H, Lou V K L, Volatility diagrams for silica,silicon-nitride, and silicon-carbide and their applicationto high-temperature decomposition and oxidation. Journal ofthe American Ceramic Society 73, 1990, 2789–2803.

134. Eriksson G, Hack K, Chemsage—A computer-program forthe calculation of complex chemical-equilibria.Metallurgical Transactions B 21, 1990, 1013–1023.

135. Narushima T, Goto T, Iguchi Y, Hirai T,High-temperature active oxidation of chemicallyvapordeposited silicon-carbide in an Ar-O 2 atmosphere.Journal of the American Ceramic Society 74, 1991,2583–2586.

136. Narushima T, Goto T, Yokoyama Y, Hagiwara J, Iguchi Y,Hirai T, High-temperature active oxidation andactive-to-passive transition of chemically vapor-depositedsilicon-nitride in N 2 -O 2 and Ar-O 2 atmospheres.Journal of the American Ceramic Society 77, 1994,2369–2375. For ca aloging purposes onl

137. Narushima T, Goto T, Hagiwara J, Iguchi Y, Hirai T,High-temperature oxidation of chemically vapordepositedsilicon-nitride in a carbon-monoxide carbon-dioxideatmosphere. Journal of the American Ceramic Society 77,1994, 2921–2925.

138. Brow R K, Pantano C G, Compositionally dependent Si 2pbinding energy shifts in silicon oxynitride thin �lms.Journal of the American Ceramic Society 69, 1986, 314–316.

139. Narushima T, Lin R Y, Iguchi Y, Hirai T, Oxidation ofchemically vapor-deposited silicon-nitride in dry oxygenat 1923 to 2003K. Journal of the American Ceramic Society76, 1993, 1047–1051.

140. Schiroky G H, Oxidation behavior of chemicallyvapor-deposited silicon-carbide. Advanced CeramicMaterials 2, 1987, 137–141.

141. Narushima T, Goto T, Yokoyama Y, Takeuchi M, Iguchi Y,Hirai T, Active-to-passive transition and bubble formation

for high-temperature oxidation of chemicallyvapor-deposited silicon-carbide in CO-CO 2 atmosphere.Journal of the American Ceramic Society 77, 1994,1079–1082.

142. Narushima T, Goto T, Hirai T, High-temperature passiveoxidation of chemically vapor-deposited silicon-carbide.Journal of the American Ceramic Society 72, 1989,1386–1390.

143. Ainslie N G, Morelock C R, Turnbull D, Symposium onNucleation and Crystallization in Glasses and Melts, M. K.Reser, G. Smith, H. Insley, Eds. The American CeramicSociety, Columbus, OH, 1962, p97.

144. Kingary W D, Introduction to Ceramics, Wiley & Sons,New York, 1960, p137.

145. Heuer A H, Ogbuji L U, Mitchell T E, Themicrostructure of oxide scales on oxidized Si and SiCsinglecrystals. Journal of the American Ceramic Society 63,1980, 354–355.

146. Costello J A, Tressler R E, Oxidation-kinetics ofhot-pressed and sintered alpha-SiC. Journal of theAmerican Ceramic Society 64, 1981, 327–331.

147. Cubicciotti D, Lau K H, Kinetics of oxidation ofhot-pressed silicon-nitride containing magnesia. Journalof the American Ceramic Society 61, 1978, 512–517.

148. Levin E M, Robbins C R, McMurdie H F, Phase Diagramfor Ceramists, 1969 Supplement, The American CeramicSociety, Columbus, OH, 1969, p88.

149. Banya S, Iguchi Y, Yamamoto S, Solubility anddissolution rate of water-vapor in liquid CaO-SiO 2 -MgOand CaO-SiO 2 -TiO 2 slags. Tetsu to Hagane—Journal of theIron and Steel Institution of Japan 72, 1986, 2210–2217.

150. Moulson A J, Roberts J P, Water in silica glass.Transactions of the Faraday Society 57, 1961, 1208–1216.

151. Meier G H, Pettit F S, The oxidation behavior ofintermetallic compounds. Materials Science and EngineeringA 153, 1992, 548–560.

152. Meier G H, Pettit F S, High-temperature oxidation andcorrosion of intermetallic compounds. Materials Scienceand Technology 8, 1992, 331–338.

153. Yao Z, Stiglich J, Sudarshan T S, Molybdenum silicidebased materials and their properties. Journal of MaterialsEngineering and Performance 8, 1999, 291–304.

154. Arvanitis A, Diplas S, Tsakiropoulos P, Watts J F,Whiting M J, Morton S A, Matthew J A D, An experimentalstudy of bonding and crystal structure modi�cations in MoSi2 and MoSi 2 + x Al (x = 10 to 40 at% Al) via Augerparameter shifts and charge transfer calculations. ActaMaterialia 49, 2001, 1063–1078.

155. Andersen O K, Jepsen O, Antonov V N, Antonov V N,Yavorsky B Y, Perlov A Y, Shpak A P, Fermisurface, bonding,and pseudogap in MoSi 2 . Physica B 204, 1995, 65–82.

156. Bartlett R W, McCamont J W, Gage P R, Structure andchemistry of oxide �lms thermally grown on molybdenumsilicides. Journal of the American Ceramic Society 48,1965, 551–558.

157. Mitra R, Rao V V R, Mahajan Y R, Oxidation behaviourof reaction hot pressed MoSi 2 -SiC composites at 500degrees C. Materials Science and Technology 13, 1997,415–419.

158. Fitzer E, Schwab J, The chemical stability ofmolybdenum disilicides as a raw material, Metal 9, 1955,1062–1066.

159. Westbrook J H, Wood D L, “PEST” degradation inberyllides, silicides, aluminides and related compounds.Journal of Nuclear Materials 12, 1964, 208–215.

160. Schlichting J, Molybdenum disilicide as a moderncomponent high temperature bonded material. HighTemperatures-High Pressures 10, 1978, 241–269.

161. Meschter P J, Low-temperature oxidation of molybdenumdisilicide. Metallurgical Transactions A 23, 1992,1763–1772.

162. Chou T C, Nieh T G, New observations of MoSi 2 pestat 500 degrees C. Scripta Metallurgica et Materialia 26,1992, 1637–1642.

163. Maloney M J, Hecht R J, Development ofcontinuous-�ber-reinforced MoSi 2 -base composites.Materials Science and Engineering A 155, 1992, 19–31.

164. Mckamey C G, Tortorelli P F, DeVan J H, Carmichael CA, A study of pest oxidation in polycrystalline MoSi 2 .Journal of Materials Research 7, 1992, 2747–2755.

165. Berztiss D A, Cerchiara R R, Gulbransen E A, Pettit FS, Meier G H, Oxidation of MoSi 2 and comparison withother silicide materials. Materials Science and EngineeringA 155, 1992, 165–181.

166. Yanagihara K, Przybylski K, Maruyama T, The role ofmicrostructure on pesting during oxidation of MoSi 2 andMo(Si, Al) 2 at 773K. Oxidation of Metals 47, 1997,277–293.

167. Berkowit J B, Blackbur P E, Felten E J,Intermediate-temperature oxidation behavior of molybdenumdisilicide. Transactions of the Metallurgical Society ofAIME 233, 1965, 1093.

168. Kurokawa K, Houzumi H, Saeki I, Takahashi H, Lowtemperature oxidation of fully dense and porous MoSi 2 .Materials Science and Engineering A 261, 1992, 292–299.

169. Chou T C, Nieh T G, Pesting of the high-temperatureintermetallic MoSi 2 . JOM 45, 1993, 15–21.

170. Yanagihara K, Maruyama T, Nagata K, Effect of thirdelements on the pesting suppression of Mo-Si-Xintermetallics(X = Al, Ta, Ti, Zr and Y). Intermetallics 4,1996, S133–S139.

171. Kingery W D, Bowen H D, Uhlmann, Introduction toCeramics, Wiley, New York, 1975, p92.

172. Fitzer E, Herbst H, Schlicting J,Schutzvonhochschmelzenden Metallenund GraphitgegenOxydationund Heißkorrosionbis1500°C durchaufgesinterteMoSi2 -Schichten. Materials and Corrosion 24, 1973, 274–282.

173. Cockeram B V, Growth and oxidation resistance ofboron-modi�ed and germanium-doped silicide diffusioncoatings formed by the halide-activated pack cementationmethod. Surface and Coatings Technology 76–77, 1995,20–27.

174. Kircher T A, Courtright E L, Engineering limitationsof MoSi 2 coatings. Materials Science and Engineering A155, 1992, 67–74.

175. Yanagihara K, Maruyama T, Nagata K, High-temperature

oxidation of Mo-Si-X intermetallics (X = Al, Ti, Ta, Zrand Y). Intermetallics 3, 1995, 243–251.

176. Yanagihara K, Maruyama T, Nagata K, Isothermal andcyclic oxidation of Mo(Si 1-x ,Al x ) 2 up to 2048K.Materials Transactions JIM 34, 1993, 1200–1206.

177. Mitra R, Rama Rao V V, Effect of minor alloying withAl on oxidation behaviour of MoSi 2 at 1200°C. MaterialsScience and Engineering A 260, 1999, 146–160.

178. Stergiou A, Tsakiropoulos P, The intermediate andhigh-temperature oxidation behaviour of (Mo, X)Si 2 (X = W,Ta) intermetallic alloys. Intermetallics 5, 1997, 117–126.

179. Cook J, Khan A, Lee E, Mahapatra R, Oxidation of MoSi2 -based composites. Materials Science and Engineering A155, 1992, 183–198.

180. Meyer M K, Akinc M, Oxidation behavior ofboron-modi�ed Mo 5 Si 3 at 800 degrees-1300 degrees C.Journal of the American Ceramic Society 79, 1996, 938–944.

181. Natesan K, Deevi S C, Oxidation behavior of molybdenumsilicides and their composites. Intermetallics 8, 2000,1147–1158.

182. Yoshimi K, Nakatani S, Suda T, Hanada S, Habazaki H,Oxidation behavior of Mo 5 SiB 2 -based alloy at elevatedtemperatures. Intermetallics 10, 2002, 407–414.

183. Park J S, Sakidja R, Perepezko J H, Coating designsfor oxidation control of Mo-Si-B alloys. ScriptaMaterialia 46, 2002, 765–770.

184. Supatarawanich V, Johnson D R, Liu C T, Oxidationbehavior of multiphase Mo-Si-B alloys. Intermetallics 12,2004, 721–725.

185. Mandal P, Thom A J, Kramer M J, Behrani V, Akinc M,Oxidation behavior of Mo-Si-B alloys in wet air. MaterialsScience and Engineering A 371, 2004, 335–342.

186. Yoshimi K, Nakatani S, Hanada S, Ko S-H, Park Y-H,Synthesis and high temperature oxidation of Mo-Si-B-Opseudo in situ composites. Science and Technology ofAdvanced Materials 3, 2002, 181–192.

187. Bewlay B P, Jackson M R, Zhao J C, Subramanian P R, Areview of very-high-temperature Ni-silicidebased

composites. Metallurgical and Materials Transactions A 34,2003, 2043–2052.

188. Rausch J J, ARF-2981-4. Armour Research FoundationNSA, 1961, 15-31171.

189. Murakami T, Sasaki S, Ichikawa K, Kitahara A,Microstructure, mechanical properties and oxidationbehavior of Nb-Si-Al and Nb-Si-N powder compacts preparedby spark plasma sintering. Intermetallics 9, 2001,621–627.

190. Zhang F, Zhang L T, Shan A D, Wu J S, In situobservations of the pest oxidation process of NbSi 2 at1023K. Scripta Materialia 53, 2005, 653–656.

191. Jackson M R, Rowe R G, Skelly D W, Oxidation of someintermetallic compounds and intermetallic matrixcomposites. Materials Research Society SymposiumProceedings 364, 1995, 1339–1344.

192. Murakami T, Sasaki S, Ichikawa K, Kitahara A,Oxidation resistance of powder compacts of the Nb-Si-Crsystem and Nb 3 Si 5 Al 2 matrix compacts prepared byspark plasma sintering. Intermetallics 9, 2001, 629–635.

193. Murakami T, Sasaki S, Ito K, Oxidation behavior andthermal stability of Cr-doped Nb(Si,Al) 2 and Nb 3 Si 5Al 2 matrix compacts prepared by spark plasma sintering.Intermetallics 11, 2001, 269–278. Fo cataloging purposesnly

194. Menon E S K, Mendiratta M G, Dimiduk D M, StructuralIntermetallics 2001, K. J. Hemker, D. M. Dimiduk, H.Clemmens, R. Darolia, H. Inui, J. M. Larsen, V. K. Sikka,M. Thomas, J. D. Whittenberger, Eds. Warrendale, PA, TMS,2001, p591.

195. Menon E S K, Mendiratta M G, High temperatureoxidation in multicomponent Nb alloys. Materials ScienceForum 475–479, 2005, 717–720.

196. Balsone S J, Bewlay B P, Jackson M R, Subramanian P R,Zhao J C, Chatterjee A, Heffernan T M, Materials beyondsuperalloys: Exploiting high-temperature composites.Structural Intemetallics 2001, 2001, 99–108.

197. Bewlay B P, Jackson M R, Zhao J-C, Subramanian P R,Mendiratta M G, Lewandowski J J, UltrahightemperatureNb-silicide-based composites. MRS Bulletin 28, 2003,

646–653.

198. Strydom W J, Lombaard J C, Pretorius R,Thermal-oxidation of the silicide CoSi 2 , CrSi 2 , NiSi 2, PtSi, TiSi 2 and ZrSi 2 . Thin Solid Films 131, 1985,215–231.

199. Kurokawa K, Yamauchi A, Classi�cation of oxidationbehavior of disilicides. Designing of interfacialstructures in advanced materials and their joints. SolidState Phenomena 127, 2007, 227–232.

200. Frampton R D, Irene E A, d’Heurle F M, A study of theoxidation of selected metal silicides. Journal of AppliedPhysics 62, 1987, 2972–2980.

201. Ye�menko L N, Nechiporenko Ye P, Pavlov V N, Fiz MetMetalloved, 16(1963) 931–937.

202. Zirinsky S, Hammer W, d’heurle F, Baglin J, Meindl JD, Oxidation mechanism in WSi 2 thin �lms. AppliedPhysical Letters 33, 1978, 76–78.

203. Mohammadi F, Saraswat K C, Meindl J D, Kinetics of thethermal-oxidation of WSi 2 . Applied Physical Letters 35,1979, 529–531.

204. Vasudevan A K, Petrovic J J, A comparative overview ofmolybdenum disilicide composites. Materials Science andEngineering A 155, 1992, 1–17.

205. Kim H-S, Yoon J-K, Kim G-H, Doh J-M, Kwun S-I, HongK-T, Growth behavior and microstructure of oxide scalesgrown on WSi 2 coatings. Intermetallics 16, 2008, 360–372.

206. Kofstad P, High Temperature Corrosion, ElsevierApplied Science, New York, 1988, p5.

207. Schwettmann F N, Graff R A, Lolodney M, Mechanism ofthe oxidation of titanium disilicide. Journal of theElectrochemical Society 118, 1971, 1973–1977.

208. Mitra R, Rao V V R, Elevated-temperature oxidationbehavior of titanium silicide and titanium silicide basedalloy and composite. Metallurgical and MaterialsTransactions A 29, 1998, 1665–1675.

209. Raj S V, An evaluation of the properties of Cr 3 Sialloyed with Mo. Materials Science and Engineering A 201,1995, 229–241.

210. Tomasi A, Ceccato R, Nazmy M, Gialanella S,Microstructure and oxidation behaviour ofchromiummolybdenum silicides. Materials Science andEngineering A 239–240, 1997, 877–881.

211. Ma J, Gu Y L, Shi L, Chen L, Yang Z, Qian Y, Synthesisand oxidation behavior of chromium silicide (Cr 3 Si)nanorods. Journal of Alloys and Compounds 375, 2004,249–252.

212. Choy K L, Chemical vapor deposition of coatings.Progress in Materials Science 48, 2003, 57–170.

213. Kung S C, Rapp R A, Kinetic-study of aluminization ofiron by using the pack cementation technique. Journal ofthe Electrochemical Society 135, 1988, 731–741.

214. Nichols E S, Burger J A, Hanink D K, Protectivecoatings for turbine parts. Mechanical Engineering 87,1965, 52–56.

215. Galmiche P M, Production of surface diffusion alloys,US Patent, 1975, No. 3,900,613.

216. Restall J E, Wood M I, Alternative processes andtreatments. Materials Science and Technology 2, 1986,225–231.

217. Parzukowski R S, Gas-phase deposition of aluminum onnickel alloys, Thin Solid Films 45, 1977, 349–355.

218. Gauje G, Morbioli R, Vapor phase aluminizing toprotect turbine airfoils, in High Temperature ProtectiveCoatings, S. C. Singhal, Ed. The Metallurgical Society ofAIME, Atlanta, GA, pp. 13–26, 1983.

219. Warnes B M, Punola D C, Clean diffusion coatings bychemical vapor deposition. Surface and Coatings Technology94–95, 1997, 1–6.

220. Levine S R, Caves R M, Thermodynamics and kinetics ofpack aluminde coating formation on IN-100. Journal of theElectrochemical Society 121, 1974, 1051–1064.

221. Goward G W, Boone D H, Mechanisms of formation ofdiffusion aluminide coatings on nickel-base superalloys.Oxidation of Metals 3, 1971, 475–495.

222. Zhang Y, Haynes J A, Lee W Y, Wright I G, Pint B A,

Cooley K M, Liaw P K, Effects of Pt incorporation on theisothermal oxidation behavior of chemical vapor depositionaluminide coatings. Metallurgical Materials Transactions A32, 2001, 1727–1741.

223. Hou P Y, McCarty K F, Surface and interfacesegregation in beta-NiAl with and without Pt addition.Scripta Materilia 54, 2006, 937–941.

224. Hou P Y, Tolpygo V K, Examination of the platinumeffect on the oxidation behavior of nickel-aluminidecoatings. Surface and Coatings Technology 202, 2007,623–627.

225. Rapp R A, Wang D, Weisert T, Simultaneouschromizing–aluminzing of iron and iron-based alloys bypack cementation, in Metallurgical Coatings, M. hobaib, R.Krutenat, Eds., TMS-AIME, Warrendale, PA, p. 131, 1987.

226. Bianco R, Rapp R A, Pack cementation aluminidecoatings on Superalloys: Codeposition of Cr and reactiveelements. Journal of the Electrochemical Society 140, 1993,1181–1190.

227. Bianco R, Rapp R A, Smialek J L, Chromium and reactiveelement modi�ed aluminide diffusion coatings onsuperalloys: Environmental testing. Journal of theElectrochemical Society 140, 1993, 1191–1203.

228. Tu D C, Lin C C, Liao S J, Chou J C, A study ofyttrium-modi�ed aluminide coatings on IN 738 alloy.Journal of Vacuum Science and Technology A 4, 1986,2601–2608.

229. Lehnert G, Meinhardt H, Present state and trend ofdevelopment of surface coating methods against oxidationand corrosion at high temperatures. Electrodeposition andSurface Treatment 1, 1972, 71–76.

230. Farrell M S, Boone D H, Streiff R, Oxide adhesion andgrowth characteristics on platinum-modi�ed aluminidecoatings. Surface and Coatings Technology 32, 1987, 69–84.

231. Göbel M, Rahmel A, Schütze M, Schorr M, Wu W T,Interdiffusion between the platinum-modi�ed aluminidecoating RT-22 and nickel-based single-crystal superalloysat 1000 and 1200°C. Materials at High Temperatures 12,1994, 301–309.

232. Warnes B M, Punola D C, Clean diffusion coatings by

chemical vapor deposition. Surface and Coatings Technology94–95, 1997, 1–6.

233. Warnes B M, Punola D C, Development of diffusioncoatings for gas turbines, in Power Generation Technology:The International Review of Preliminary Power Production,R. Knox, Ed., Sterling Publications Limited, London, pp.207–210, 1995.

234. Yu R, Hou P Y, First-principles calculations of theeffect of Pt on NiAl surface energy and the site preferenceof Pt. Applied Physics Letters 91, 2007, 1−3.

235. Wu W T, Rahmel A, Schorr M, Role of platinum in the Na2 SO 4 -induced hot corrosion resistance of aluminumdiffusion coatings. Oxidation of Metals 22, 1984, 59–81.

236. Schaeffer J, Kim G M, Meier G H, Pettit F S, Theeffects of precious metals on the oxidation and hotcorrosion of coatings, in The Role of Active Elements inthe Oxidation Behavior of High Temperature Metals andAlloys, Lang E, Ed., Elsevier, Amsterdam, pp. 231–267,1989.

237. Leyens C, Pint B A, Wright I G, Effect of compositionon the oxidation and hot corrosion resistance of NiAldoped with precious metals. Surface and Coatings Technology133, 2000, 15–22.

238. Smialek J L, Oxide morphology and spalling model forNiAl. Metallurgical Transactions A 9, 1978, 309–320.

239. Christensen R J, Lipkin D M, Clarke D R, The stressand spalling behavior of the oxide scale formed onpolycrystalline Ni 3 Al. Acta Materialia 44, 1996,3813–3821.

240. Pérez P, González-Carrasco J L, Adeva P, Oxidationbehavior of a Ni 3 Al PM alloy. Oxidation of Metals 48,1997, 143–170.

241. Haanappel V A C, Pérez P, González-Carraco J L,Stroosnijder M F, The cyclic oxidation behaviour of PM Ni3 Al in air at elevated temperatures: The effect of Cr- andY-implantation. Intermetallics 6, 1998, 347–356.

242. Kumar A, Nasralla M M, Douglass D L, Effect of yttriumand thorium on oxidation behavior of Ni-Cr-Al alloys.Oxidation of Metals 8, 1974, 227–263.

243. Shankar S, Seigle L L, Interdiffusion and intrinsicdiffusion in NiAl(delta) phase of Al-Ni system.Metallurgical Transactions A 9, 1978, 1467–1476.

244. Xu C, Peng X, Wang F, Cyclic oxidation of anultra�ne-grained and CeO 2 -dispersed δ-Ni 2 Al 3 coating.Corrosion Science 52, 2010, 740–747.

245. Xu C, Peng X, Zheng L, Wang F, Erosion-corrosion in alaboratory coal-�ring FBC of various aluminized coatingsprepared by low-temperature pack cementation. Surface andCoatings Technology 205, 2011, 4540–4546.

246. Zhan Z L, He Y D, Wang D R, Gao W, Low-temperatureprocessing of Fe-Al intermetallic coatings assisted byball milling. Intermetallics 14, 2006, 75–81.

247. Samuel R L, Lockidngton N A, The protection ofmetallic surfaces by chromium diffusion–Application ofChromizing: Part I-VII, Metal Treatment and Drop Gorging,18, 1951, 354–359, 407–414, 440–444, 495–506, 543–548, and19, 1952, 27–32, 81–85.

248. Drewett R, Diffusion coatings for the protection ofiron and steel. Anti-Corrosion Methods and Materials 16,1969, 11–16.

249. Peng X, Yan J, Xu C, Wang F, Oxidation at 900°C of thechromized coatings on A3 carbon steel with theelectrodeposition pretreatment of Ni or Ni-CeO 2 �lm.Metallurgical and Materials Transactions A 39, 2008,119–129. F r catalogi g purpos s o ly

250. Yan J, Peng X, Wang F, Oxidation of a novel CeO 2-dispersion-strengthened chromium coating in simulatedcoal-combustion gases. Materials Science and Engineering A426, 2006, 266–273.

251. Zhu L, Peng X, Yan J, Wang F, Wu W, Oxidation of anovel chromium coating with CeO 2 dispersions. Oxidationof Metals 62, 2004, 411–426.

252. Maziasz J, Swindeman R W, Selecting and developingadvanced alloys for creep-resistance for microturbinerecuperator applications. Journal of Engineering for GasTurbines and Power 125, 2003, 310–315.

253. Saunders S R J, Nicholls J R, Coatings and surfacetreatments for high temperature oxidation resistance.Materials Science and Technology 5, 1989, 780–798.

254. Meier G H, Cheng C, Perkins RA, Bakker W, Diffusionchromizing of ferrous alloys. Surface and CoatingsTechnology 39, 1989, 53–64.

255. Zhang H, Peng X, Zhao J, Wang F, Priorelectrodeposition of nanocrystalline Ni-CeO 2 �lmfabricating an oxidation-resistant chromized coating oncarbon steels. Electrochemical and Solid State Letters 10,2007, C12–C15.

256. Geib F D, Rapp R A, Simultaneouschromizing-aluminizing coating of low-alloy steels by ahalideactivated, pack-cementation process. Oxidation ofMetals 40, 1993, 213–228.

257. Vojtech D, Kubatik T, Jurek K, Maixner J,Cyclic-oxidation resistance of protective silicide layerson titanium. Oxidation of Metals 63, 2005, 305–323.

258. Li X Y, Taniguchi S, Matsunaga Y, Nakagawa K, FujitaK, In�uence of siliconizing on the oxidation behavior of agamma-TiAl based alloy. Intermetallics 11, 2003, 143–150.

259. Xiong H P, Mao W, Xie Y H, Ma W L, Chen Y F, Li X H,Li J P, Cheng Y Y, Liquid-phase siliconizing by Al-Sialloys at the surface of a TiAl-based alloy and improvementin oxidation resistance. Acta Materialia 52, 2004,2605–2620.

260. Liang W, Ma X X, Zhao X G, Zhang F, Shi J Y, Zhang J,Oxidation kinetics of the pack siliconized TiAlbased alloyand microstructure evolution of the coating. Intermetallics15, 2007, 1–8.

261. Kubatik T, Jaglova M, Kalabisova E, Cihal V,Improvement of oxidation resistance of TiAl 6 V 4 alloy bysiliconizing from liquid phase using melts with highsilicon content. Journal of Alloys and Compounds 509,2011, 5493–5499.

262. Koo C H, Yu T H, Pack cementation coatings on Ti 3Al-Nb alloys to modify the high-temperature oxidationproperties. Surface and Coatings Technology 126, 2000,171–180.

263. Suilik S B A, Takeshita K, Kitagawa H, Tetsui T,Hasezaki K, Preparation and high temperature oxidationbehavior of refractory disilicide coatings for gamma-TiAlintermetallic compounds. Intermetallics 15, 2007,

1084–1090.

264. Tsirlin M S, Kasatkin A V, Byalobzheskii A V, Anoxidation-resistant silicide coating for niobium alloys.Soviet Powder Metallurgy and Metal Ceramics 17, 1978,920–923.

265. Suzuki R O, Ishikawa M, Ono K, MoSi 2 coating onmolybdenum using molten salt. Journal of Alloys andCompounds 306, 2000, 285–291.

266. Talboom F P, Grafwallner J, Nickel or cobalt base witha coating containing iron chromium and aluminum, US Patent,1970, No 3,542,530.

267. Mom A J A, High Temperature Coatigns for Gas Trubine,an Overiew, NLR Report, MP 81003U, Amsterdam, 1981.

268. Li C J, Li W Y, Effect of sprayed powder particle sizeon the oxidation behavior of MCrAlY materials during highvelocity oxygen-fuel deposition. Surface and CoatingsTechnology 162, 2003, 31–41.

269. Evans D J, Elam R C, Cobalt base coating for thesuperalloys, US Patent, 1972, No. 3,676,085.

270. Hecht R J, Goward G W, Elam R C, High temperatureNiCoCrAlY coating, US Patent, 1975, No. 3,928,026.

271. Gil A, Shemet V, Vassen R, Subanovic M, Toscano J,Naumenko D, Singheiser L, Quadakkers W J, Effect ofsurface condition on the oxidation behavior of MCrAlYcoatings. Surface and Coatings Technology 201, 2006,3824–3828.

272. Yanar N M, Meier E M, Pettit F S, Meier G H, Theeffect of yttrium content on the oxidationinduceddegradation of NiCoCrAlY coatings and the in�uence ofsurface roughness on the oxidation of NiCoCrAlY andplatinum aluminide coatings. Materials at High Temperatures26, 2009, 331–338.

273. Wang F H, Lou H Y, Wu W T, In�uence of defects onproperties of sputtered CoCrAlY coating. Acta MetallurgicaSinica 28, 1992, B443–B448.

274. Taniguchi S, Asanuma N, Shibata T, Wang F H, Lou H Y,Wu W T, Oxidation resistance of TiAl improved by CoCrAland CoCrAlY coating. Journal of Japan Institute of Metals57, 1993, 781–789.

275. Choquet P, Indrigo C, Mevrel R, Microstructure ofoxide scales formed on cyclically oxidized M-CrAl-Ycoatings. Materials Science and Engineering 88, 1987,97–101.

276. Brandla W, Toma D, Grabke H J, The characteristics ofalumina scales formed on HVOF-sprayed MCrAlY coatings.Surface and Coatings Technology 108, 1998, 10–15.

277. Strauss D, Muller G, Schumacher G, Engelko V, Stamm W,Clemens D, Quaddakers W J, Oxide scale growth on MCrAlYbond coatings after pulsed electron beam treatment anddeposition of EBPVD-TBC. Surface and Coatings Technology135, 2001, 196–201.

278. Frances M, Vilasi M, Mansourgabr M, Steinmetz J,Steinmetz P, Study of the oxidation at 850°C of the alloyNi-Co-Cr-Al(-Y)-Ta cast and sprayed by a low-pressureplasma spraying process. Materials Science and Engineering88, 1987, 89–96.

279. Haynes J A, Ferber M K, Porter W D, Damageaccumulation in thermally-cycled Al 2 O 3 scales on plasmasprayed NiCrAlY, in Proceedings of the Symposium on HighTemperature Corrosion and Materials Chemistry, San Diego,CA, 1998, pp. 146–157.

280. Ajdelsztajn L, Picas J A, Kim G E, Bastian F L,Schoenung J, Provenzano V, Oxidation behavior of HVOFsprayed nanocrystalline NiCrAlY powder. Materials Scienceand Engineering A 338, 2002, 33–43.

281. Gupta D K, Metallurgical characterization ofplasma-sprayed M-Cr-Al-Y coatings. Thin Solid Films 64,1979, 335–335.

282. Wang F, Tian X, Li Q, Li L, Peng X, Oxidation and hotcorrosion behavior of sputtered nanocrystalline coating ofsuperalloy K52. Thin Solid Films 516, 2008, 5740–5747.

283. Chen G F, Lou H Y, Predicting the oxide formation ofNi-Cr-Al alloys with nano-sized grain. Materialia Letters45, 2000, 286–291.

284. Yang X, Peng X, Xu C, Wang F, Electrochemical assemblyof Ni-xCr-yAl nanocomposites with excellenthigh-temperature oxidation resistance. Journal of theElectrochemical Soceity 156, 2009, C167–175.

285. Li Q, Peng X, Zhang J Q, Zong G X, Wang F H,Comparison of the oxidation of high-sulfur Ni-25Cr-5Alalloys in as-cast and as-sputtered states. CorrosionScience 52, 2010, 1213–1221.

286. Yang S, Wang F, Effect of nanocrystallization onsulfur segregation in Fe-Cr-Al alloy during oxidation at1000°C. Oxidation of Metals 65, 2006, 195–205.

287. Klam H J, Hahn H, Gleiter H, The thermal expansion ofgrain boundaries. Acta Materialia 35, 1978, 2101–2104.

288. Ma W, Dong H, Ceramic thermal barrier coatingmaterials, in Thermal Barrier Coating, H. Xu, H. Guo, Eds.Woodhead Publishing Limited, Cambridge, U.K., pp. 25–44,2011.

289. Jones R L, Development of hot-corrosion-resistantzirconia thermal barriers coatings. Materials at HighTemperatures 9, 1991, 228–236.

290. Jones R L, Some aspects of the hot corrosion of thethermal barrier coatings. Journal of Thermal SprayTechnology 6, 1997, 77–84.

291. Cao X Q, Vassen R, Stoever D, Ceramic materials forthermal barrier coatings. Journal of the European CeramicSociety 24, 2004, 1–10.

292. Cao X, Vassen R, Fischer W, Tietz F, Jungen W, StoeverD, Lanthanum-cerium oxide thermal barrier coating materialfor high-temperature applications. Advanced Materials 15,2003, 1438–1442.

293. Ma W, Gong S, Xu H, Cao X, The thermal cyclingbehavior of lanthanum-cerium oxide thermal barrier coatingprepared by EB-PVD. Surface and Coatings Technology 200,2006, 5113–5118.

294. Zhou C G, Yu Q H, Nanostructured thermal barriercoatings, in Thermal Barrier Coating, H Xu and H Guo, Eds.Woodhead Publishing Limited, Cambridge, U.K., pp. 75–96,2011.

295. Tolpygo V K, Clarke D R, Surface rumpling of a (Ni,Pt)Al bond coat induced by cyclic oxidation. ActaMaterialia 48, 2000, 3283–3293.

296. Tolpygo V K, Clarke D R, On the rumpling mechanism innickel-aluminide coatings: Part I: An experimental

assessment. Acta Materialia 52, 2004, 5115–5127.

297. Tolpygo V K, Clarke D R, Rumpling of CVD (Ni,Pt)Aldiffusion coatings under intermediate temperature cycling.Surface and Coatings Technology 203, 2009, 3278–3285.

298. Evans A G, Mumm D R, Hutchinson J W, Meier G H, PettitF S. Mechanisms controlling the durability of thermalbarrier coatings. Progress in Materials Science 46, 2001,505–553.

299. Cao F, Tryon B, Torbet C J, Pollock T M,Microstructural evolution and failure characteristics of aNiCoCrAlY bond coat in “hot spot” cyclic oxidation. ActaMaterialia 57, 2009, 3885–3894.

300. Sergo V, Clarke D R, Observation of subcritical spallpropagation of a thermal barrier coating. Journal of theAmerican Ceramic Society 81, 1998, 3237–3242.

301. Hutchinson J W, Suo Z, Mixed-mode cracking in layeredmaterials. Advances Applied Mechanics 29, 1992, 63–191.

302. Evans A G, Cannon R M, Stresses in oxide �lms andrelationships with cracking and spalling. MaterialsScience Forum 43, 1989, 243–268. For cataloging urposesonly

6 Chapter 6 - Thermal Spray Coatings forAeronautical and Aerospace Applications

Akedo J. 2008. Room temperature impact consolidation (RTIC)of �ne ceramic powder by aerosol deposition method andapplications to microdevices. J. Therm. Spray Technol.17:181–198.

Agarwal A., Seal S., and McKechnie T. 2002. Net shapenanostructured aluminum oxide structures fabricated byplasma spray forming. J. Therm. Spray Technol. 12:350–359.

Alkimov A.P., Kosarev V.F., and Papyrin A.N. 1990. A methodof cold gas dynamic deposition. Dokl. Akad. Nauk SSSR315:1062–1065.

An K. 1999. Thermal conductivity of various thermal barriercoatings and its rationalization on the basis of theeffects of microstructure. PhD dissertation. The Universityof Utah.

Arata Y. 1988. Advanced new technology for thermalspraying. Proceedings of International Symposium onAdvanced Thermal Spraying Technology and Allied Coatings(ATTAC’88), Japan High Temperature Society, Osaka, 1988,pp. 1–8.

Arata Y., Ohmori A., and Li C.-J. 1987. Characteristics ofmetal electroplating to plasma sprayed ceramic coatings.Trans. Jpn. Weld. Res. Inst. 16:259–265.

Arata Y., Ohmori A., and Li C.-J. 1988a. Study on thestructure of plasma sprayed ceramic coating by using copperelectroplating. Proceedings of International Symposium onAdvanced Thermal Spraying Technology and Allied Coatings(ATTAC’88), Japan High Temperature Society, Osaka, 1988,pp. 205–210.

Arata Y., Ohmori A., and Li C.-J. 1988b. Electrochemicalmethod to evaluate the connected porosity in ceramiccoatings. Thin Solid Films 156:315–325. For cat logingpurp ses only

ASTM C177. 1990. Test Method for Steady-State ThermalTransmission Properties by Means of the Guarded Hot Plate.ASTM Standards, American Society for Testing and Materials,Philadelphia.

ASTM E1225. 2004. Standard Test Method for ThermalConductivity of Solids by Means of the

GuardedComparative-Longitudinal Heat Flow Technique. ASTMStandards, American Society for Testing and Materials,Philadelphia.

ASTM E1461. 2002. Standard Test Method for ThermalDiffusivity by the Flash Method. ASTM Standards, AmericanSociety for Testing and Materials, Philadelphia.

Ault N.N. 1957. Characteristics of refractory oxidecoatings produced by �ame-spraying. J. Am. Ceram. Soc.40:69–74.

Balani K., Agarwal A., and McKechnie T. 2006. Near netshape fabrication via vacuum plasma spray forming. Trans.Indian Institute Metals 59:237–244.

Barbezat G., and Landes K. 2000. Plasma technology TRIPLEXfor the deposition of ceramic coatings in industry. InThermal Spray: Surface Engineering via Applied Research,Proceedings of the 1st International Thermal SprayConference, (ed.) C.C. Berndt, ASM International Ohio, pp.881–885.

Barbezat G., Muller E., and Walser B. 1988. Applyingtungsten carbide cobalt coatings by high velocitycombustion spraying. Sulzer Tech. Rev. 4:27–33.

Bardi U., Giolli C., Scrivani A. et al. 2008. Developmentand investigation on new composite and ceramic coatings aspossible abradable seals. J. Therm. Spray Technol.17:805–811.

Bejan A. and Kraus A.D. 2003. Heat Transfer Handbook. JohnWiley & Sons, Hoboken New Jersey, USA, p. 2.

Berndt C.C., and McPherson R. 1980. A fracture mechanicsapproach to the adhesion of �ame and plasma sprayedcoatings. Proc. 9th Int. Thermal Spraying Conf, Hague,1980, Nederland Inst. Voor Lastetechnik, 1980, pp.310–316.

Baik K.H., and Grant P.S. 1999. Microstructural evaluationof monolithic and continuous �bre reinforced Al-12wt.%Siproduced by low pressure plasma spraying. Mater. Sci. Eng.A 265:77–86.

Bi X., Xu H., and Gong S. 2000. Investigation of thefailure mechanism of thermal barrier coatings prepared byelectron beam physical vapor deposition. Surf. Coat.Technol. 130:122–27.

Binder K., Gartner F., and Klassen T. 2010. Ti-parts foraviation industry produced by cold spraying. In ThermalSpray: Global Solutions for Future Application, DVS, ASMInternational and IIW, pp. 587–592.

Bisson J., Moreau C., Dorfman M., Dambra C., and Mallon J.2005. In�uence of hydrogen on the microstructure ofplasma-sprayed yttria-stabilized zirconia coatings. J.Therm. Spray Technol. 14:85–90.

Boire-Lavigne S., Moreau C., and Saint-Jacques R.G. 1995.The relationship between the microstructure and thermaldiffusivity of plasma-sprayed tungsten coatings. J. Therm.Spray Technol. 4:261–267.

Bolot R., Antou G., Montavon G., and Coddet C. 2005. Atwo-dimensional heat transfer model for thermal barriercoating average thermal conductivity computation. NumericalHeat Transfer Part A 47:875–898.

Borel M.O., Nicoll A.R., Schapfer H.W., and Schmid R.K.1989. The wear mechanisms occurring in abradable seals ofgas turbines. Surf. Coat. Technol. 39/40:117–126.

Boulos M.I., Fuachais P., Vardelle A., and Pfender E. 1993.Fundamentals of plasma particle momentum and heattransfer. In Plasma Spraying: Theory and Applications,(ed.) R. Suryanarayanan, World Scienti�c Publishing Co.,Singapore, 1993.

Bourdin E., Fauchais P., Vardelle A., and Pfender E.1983.Transient heat conduction under plasma conditions. Int. J.Heat Mass Transfer 26:567–582.

Bowden F.P., and Tabor D. 1964. Friction and Lubrication ofSolids. Oxford Press, London, UK, p. 25.

Cao X.Q. 2007. Thermal Barrier Coating Materials. SciencePublication, Beijing (in Chinese).

Cao X.Q., Vassen R., and Stoever D. 2004, Ceramic materialsfor thermal barrier coatings. J. Eur. Ceram. Soc. 24:1–10.

Chen W.R., Wu X., Marple B.R., Nagy D.R., and Patnaik P.C.2008. TGO growth behaviour in TBCs with APS and HVOF bondcoats. Surf. Coat. Technol. 202:267–2683.

Chen X., Zhao Y., Fan X., Liu Y., Zou B., Wang Y., Ma H.,and Cao X. 2011. Thermal cycling failure of new LaMgAl 11

O 19 /YSZ double ceramic top coat thermal barrier coatingsystems. Surf. Coat. Technol. 205:3293–3300.

Chivavibul P., Watanabe M., Kuroda S., and Shinoda K. 2007.Effects of carbide size and Co content on themicrostructure and mechanical properties of HVOF-sprayedWC-Co coatings. Surf. Coat. Technol. 202:509–521.

Clegg M.A., and Mehta M.H. 1988. NiCrAl/bentonite thermalspray powder for high temperature abradable seals. Surf.Coat. Technol. 34:69–77.

Clyne T.W., Golosnoy I.O., Tan J.C., and Markaki A.E. 2006.Porous materials for thermal management under extremeconditions, Phil. Trans. R. Soc. A 364:125–146.

Curry N., Markocsan N., Li X.H., Tricoire A., and DorfmanM. 2011. Next generation thermal barrier coatings for thegas turbine industry. J. Therm. Spray Technol. 20:108–115.

Davis J.R. 1997. Protective Coating for Superalloys, ASMSpecialty Handbook Heat-Resistant Materials, ASMInternational, Materials Park, Ohio, pp. 335–344.

Davis J.R. 2004. Handbook of Thermal Spray Technology, ASMInternational, Materials Park, Ohio.

Diltemiz S.F. 2010. Optimization of Thermal and MechanicalProperties of Thermal Barrier Coatings, PhD dissertation,Eskisehir Osmangazi University.

Dorfman M., Erning U., and Mallon J. 2002. Gas turbines useabradable coatings for clearance-control seals. Seal.Technol. 2002:7–8.

Dresvin S.V. 1972. Physics and Technique of Low TemperaturePlasma, Atomizdat, Moskva, Russia (in Russian).

Eshelby J.D. 1959. The elastic �eld outside an ellipsoidalinclusion, Proc. Roy. Soc. London Series-A. 252:562–569.

Espie G., Fauchais P., Labbe J.C., Vardelle A., andHannoyer B. 2001. Oxidation of iron particles during APS:Effect of the process on formed oxide, wetting of dropletson ceramics substrates. In Thermal Spray 2001: NewSurfaces for a New Millennium (eds.) C.C. Berndt, K.A.Khor, and E. Lugscheider E. ASM International, Ohio, pp.821–827.

Evans A.G., Mumm D.R., Hutchinson J.W., Meier G.H., and

Petitt F.S. 2001. Mechanisms controlling the durability ofthermal barrier coatings. Progress Mater. Sci. 46:505–553.

Exner H. E., and Gurland J. 1970. A review of parametersin�uencing some of mechanical properties of tungstencarbide-cobalt alloys. Powder Metall. 13:13–63.

Fan S.-Q., Yang G.-J., Li C.-J., Liu G.-J., Li C.-X., andZhang L.-Z. 2006. Characterization of microstructure ofTiO 2 coating deposited by vacuum cold spraying. J. Therm.Spray Technol. 15:513–517.

Fang J.C., and Xu W.J. 2002. Plasma spray forming. J.Mater. Process. Technol. 129, 288–293.

Faraoun H.I., Grosdidier T., Seichepine J.L. et al. 2006.Improvement of thermally sprayed abradable coating bymicrostructure control. Surf. Coat. Technol. 201:2303–2312.

Fauchais P., Vardelle M., Vardelle A., and Coudert J.F.1989. Plasma spraying of ceramic particles in argonhydrogenD.C. plasma jets: Modeling and measurements of particles in�ight correlation with thermophysical properties of sprayedlayer. Metall. Trans. 20B:263–276.

Franck J., and Kingery W.D. 1954. Thermal conductivity: IX,Experimental investigation of effect of porosity onthermal conductivity. J. Amer. Ceram. Soc. 37:99–107.

Fukanuma H., and Ohno N. 2003. In�uence of substrateroughness and temperature on adhesive strength in thermalspray coatings. In Thermal Spray 2003, Advancing theScience & Applying the Technology (ed.) C. Moreau and B.Marple, ASM International, Materials Park, Ohio, pp.1361–1368.

Gorham 1990. The Expending Business Opportunities andChallenges in Thermal Spray (TS) Coatings. Gorham AdvancedMaterials Institute, London, UK, Vol. 1.

Geçkinli A.E. 1991. İleri Teknoloji Malzemeleri. İstanbulTechnical University, Istanbul, Turkey, p. 250.

Gerling R., Schimansky F.P., Wegmann G., and Zhang J.X.2002. Spray forming of Ti48.9Al (at.%) and subsequent hotisostatic pressing and forging. Mater. Sci. Eng. A326:73–78.

Golosnoy I.O., Tsipas S.A., and Clyne T.W. 2005. Ananalitical modeling for simulation of heat �ow in

plasmasprayed thermal barrier coatings. J. Therm. SprayTechnol. 14:205–214.

Guo H.B., Kuroda S., and Murakami H. 2006. Segmentedthermal barrier coatings produced by atmospheric plasmaspraying hollow powders. Thin Solid Films 506–507:136–139.

Guo H.B., Vaßen R., and Stöver D. 2004. Atmospheric plasmasprayed thick thermal barrier coatings with highsegmentation crack density. Surf. Coat. Technol.186:353–363.

Gyenis L., Grimaud A., Betoule O., Monerie-Moulin F., andFauchais P. 1991. In�uence of temperature control duringspraying on hardness and cohesion of alumina coatings. InProceedings of 2nd PlasmaTechnik-Symposium (eds.) S.Blum-Sandmeier, H. Eschnauer, P. Huber, and A.R. Nicoll,Ha�iger Druck AG, Wettingen, Vol. 1, pp. 95–101.

Hao S., Li C.-J., and Yang G.-J. 2011. In�uence ofdeposition temperature on the microstructures andproperties of plasma-sprayed Al 2 O 3 coatings. J. Therm.Spray Technol. 20:160–169.

Hawthorne H.M., Erickson L.C., Ross D., Tai H., andTroczynski T. 1997. The microstructural dependence of wearand indentation behavior of some plasma-sprayed aluminacoatings. Wear 203–204:709–714.

Hedges M.K., Newbery A.P., and Grant P.S. 2002.Characterisation of electric arc spray formed Ni superalloyIN718. Mater. Sci. Eng. A 326:79–91.

Houben J.M., and Liempd G.G. 1983. Metallurgicalinteractions of Mo and steel during plasma spraying.Proceedings of 10th International Thermal Spray Conference,DVS-German Welding Research Institute, Dusseldolf, pp.

Huang R.Z., Wang Y.Y., Yang G.J., Li C.J., and Li Q.L.2004. Effect of hot pressing pressure on the microstructureand mechanical property of plasma sprayed B/Al composite.Aerospace Mater. Technol. 2004:51–55 (in Chinese). Forcataloging pu p ses only

Ilavsky J., Allen A.J., Berndt C., and Herman H. 1995.Anisotropy of the surfaces of pores in plasma sprayedalumina deposits. In Thermal Spraying: Current Status andFuture Trends (Proceedings of 14th ITSC) (ed.) A. Ohmori,Japan High Temperature Society, Osaka, pp. 483–488.

Incropera F.P., and DeWitt D.P. 2007. Fundamentals of Heatand Mass Transfer. John Wiley & Sons, Canada.

Iwamoto B. and Umesaki N. 1985. Characterization of plasmasprayed zirconia coatings by x-ray diffraction and Ramanspectroscopy. Thin Solid Films 127:129–137.

Jacobs L., Hyland M.M., and De Bonte M. 1999. Study of thein�uence of microstructural properties on the sliding-wearbehavior of HVOF and HVAF sprayed WC-cermet coatings. J.Therm. Spray Technol. 8:125–132.

Ji G.-C., Li C.-J., and Wang Y.-Y. 2006. Microstructuralcharacterization and abrasive wear performance of HVOFsprayed Cr3C2-NiCr coating. Surf. Coat. Technol.200:6749–6757.

Jiang X., Liu C., and Lin F. 2007. Overview on thedevelopment of nanostructured thermal barrier coatings.J. Mater. Sci. Technol. 23:449–456.

Jones H. 1971. Cooling, freezing and substrate impact ofdroplets formed by rotary atomization. J. Phys. D: Appl.Phys. 4:1657–1660.

Jones H. 1973. Splat cooling and metastable phases. ReportProgress Phy. 36:425–1497.

Kawase R. 1993. Evaluation methods for thermal spraycoatings. Proceedings of the 1993 National Thermal SprayConference, ASM International, Ohio, Anaheim, USA, pp.575–579.

Kawase R., Nakajima H., Maehira K., Koyama M., and Abe H.1988. In�uence of spraying parameters on detonationcoating properties. Proceedings of International Symposiumon Advanced Thermal Spraying Technology and AlliedCoatings (ATTAC’88), Japan High Temperature Society, Osaka,1988, pp. 55–60.

Kim H.-J., Lee C.-H., and Hwang S.-Y. 2005. Superhard nanoWC–12%Co coating by cold spray deposition. Mater. Sci.Eng. A 391:243–248.

Kobayashi A., and Arata Y. 1988. Formation of high qualityceramic coatings by gas tunnel type plasma spraying.Proceedings of International Symposium on Advanced ThermalSpraying Technology and Allied Coatings (ATTAC’88), JapanHigh Temperature Society, Osaka, 1988, pp. 107–112.

Kowalsky K.A., Marantz D.R., Smith. M.F., and OberkamptW.L. 1990. HVOF: Particle, �ame diagnostics and coatingcharacteristics. In Thermal Spray Research, Proceedings ofthe National Thermal Spray Conference, ASM International,Ohio, Long Beach, USA, pp. 877–592.

Kreye H., Schwetzke R., and Zimmermann S. 1996. Highvelocity oxy-fuel �ame spraying—Process and coatingcharacteristics. In Thermal Spray: Practical Solutions forEngineering Problems (ed.) C.C. Berndt, ASM International,Materials Park, Ohio, USA, pp. 451–456.

Kulkarni A., Gutleber J., Sampath S., Goland A., LindquistW.B., Herman H., Allen A.J., and Dowd B. 2004. Studies ofthe microstructure and properties of dense ceramic coatingsproduced by high-velocity oxygen-fuel combustion spraying.Mater. Sci. Eng. A 369:124–137.

Kuroda S. 1995. Evaluation of the pore structure inplasma-sprayed coatings. In Advances in Inorganic Filmsand Coatings (ed.) P. Vincenzini, Techna Srl., Florence,Italy, pp. 373–380.

Kuroda S., and Clyne T.W. 1991. The quenching stress inthermally sprayed coatings. Thin Solid Films 200:49–66.

Kuroda S., Kawakita J., Watanabe M., and Katanoda H. 2008.Warm spraying—A novel coating process based onhigh-velocity impact of solid particles. Sci. Technol. Adv.Mater. 9(3)Article No. 033002.

Karthikeyan J., Sinha A.K., and Biswas A.R. 1996.Impregnation of thermally sprayed coatings formicrostructural studies. J. Therm. Spray Technol. 5, 74–78.

Lawn R.B., and Wilshaw T.W. 1977. Fracture of BrittleMaterials. Cambridge University Press, New York, p. 65.

Lewis J.A., and Gauvin W.H. 1973. Motion of particlesentrained in a plasma jet. AICheE J. 19:982–990.

Leyens C., Schulz U., Pint B.A., and Wright I.G. 1999.In�uence of electron beam physical vapor deposited thermalbarrier coating microstructure on thermal barrier coatingsystem performance under cyclic oxidation conditions, Surf.Coat. Technol. 120–121:68–76.

Li C.-J. 2010. Thermal spraying of light alloys. In SurfaceEngineering of Light Alloys (ed.) H.S. Dong. WoodheadPublishing Limited, Cambridge, UK, p. 185.

Li C.-J., Ji G.-C., Wang Y.-Y., and Sonoya K. 2002. Thedominant mechanism of carbon loss during HVOF spraying ofCr 3 C 2 -NiCr. Thin Solid Films 419:137–143.

Li L., Kharas B., Zhang H., and Sampath S. 2007.Suppression of crystallization during high velocity impactquenching of alumina droplets: Observations andcharacterization. Mater. Sci. Eng. A 456:35–42.

Li C.-J, and Li W.-Y. 2003. Effect of sprayed powderparticle size on the oxidation behaviour of MCrAlYmaterials during HVOF deposition. Surf. Coat. Technol.162(1):31–41.

Li C.-J., and Li J.-L. 2004. Evaporated-gas-inducedsplashing model for splat formation during plasma spraying.Surf. Coat. Technol. 184:13–23.

Li C.-J., and Li W.-Y. 2009. Recent advances in coatingdevelopment by cold spraying. In Surface Modi�cation ofMaterials by Coatings and Films (eds.) X. Liu, P.K. Chu,and C. Ding, Research Signpost, Kerala, India, pp. 1–107.

Li C.-J., Li W.-Y., and Liao H. 2006a. Examination of thecritical velocity for deposition of particles in coldspraying. J. Therm. Technol. 15:212–222.

Li C.-J., Li C.-X., and Wang M. 2005b. Effect of sprayparameters on the electrical conductivity of plasmasprayedLa 1−x Sr x MnO 3 coating for the cathode of SOFCs. Surf.Coat. Technol. 198:278–282.

Li C.-J., Li J.-L., Wang W.-B., Ohmori A., and Tani K.1998. Effect of particle substrate materials combinationson morphology of plasma sprayed splats. In Proceedings of15th International Thermal Spray Conference (15th ITSC),May 25–29, 1998, Nice, France (ed.) C. Coddet, ASMInternational, Materials Park, USA, pp. 481–488.

Li C.-J., Li Y., Yang G.-J., and Li C.-X. 2012. A novelplasma-sprayed durable thermal barrier coating with awell-bonded YSZ interlayer between porous YSZ and bondcoat. J. Thermal Spray Technol. 21:383–390.

Li C.-J., Li C.-X., Yang G.-J., and Wang Y.-Y. 2006c.Examination of substrate surface melting-induced splashingduring splat formation in plasma spraying. J. Therm. SprayTechnol. 15:717–724.

Li Y., Li C.-J., Yang G.-J., and Xing L.-K. 2010b. Thermalfatigue behavior of thermal barrier coatings with theMCrAlY bond coats by cold spraying and low-pressure plasmaspraying. Surf. Coat. Technol. 205:2225–2233.

Li Y., Li C.-J., Zhang Q., Yang G.-J., and Li C.-X. 2010a.In�uence of TGO composition on the thermal shock lifetimeof thermal barrier coatings with cold-sprayed MCrAlY bondcoat. J. Therm. Spray Technol. 19:168–177.

Li Y., Li C.-J., Zhang Q., Yang G.-J., and Li C.-X. 2011.In�uence of TGO composition on the thermal shock lifetimeof thermal barrier coatings with cold-sprayed MCrAlTaY bondcoat. J. Therm. Spray Technol. 20:121–131.

Li C.-J., Liao H.-L., Gougen P., Montovon G., and Coddet C.2005a. Experimental determination of Relationship between�attening degree and Reynolds numbers for molten spraydroplet. Surf. Coat. Technol. 191:375–383.

Li C.-J., Ning X.-J., and Li C.-X. 2005c. Effect ofdensi�cation process on the properties of plasma-sprayedYSZ electrolyte coatings for solid oxide fuel cell. Surf.Coat. Technol. 190:60–64.

Li C.-J., and Ohmori A. 1996. The lamellar structure of adetonation gun sprayed Al 2 O 3 coating. Surf. Coat.Technol. 82:254–258.

Li C.-J., and Ohmori A. 2002. Relationship between thestructure and properties of thermally sprayed deposits. J.Therm. Spray Technol. 11:365–374.

Li C.-J., Ohmori A., and Arata Y. 1995a. Effect of spraymethods on the lamellar structure of Al 2 O 3 coatings.In Thermal Spray, Current Status and Future Trends (ed.) A.Ohmori, Japan High Temperature Society, Osaka, Japan, pp.501–506.

Li C.-J., Ohmori A., and Arata Y. 1995c. Evaluation oflamellar bonding of ceramic coatings by particle erosiontest, Thermal Spray, Current Status and Future Trends, A.Ohmori (ed.), Japan High Temperature Society, Osaka,Japan, 967–972.

Li C.-J., Ohmori A., and Harada Y. 1995b. Effect ofWC-particle size on the formation of HVOF sprayed WC-Cocoatings. In Thermal Spray, Current Status and FutureTrends (ed.) A. Ohmori, Japan High Temperature Society,Osaka, Japan, pp. 869–875.

Li C.-J., Ohmori A., and Arata Y. 1995c. Evaluation oflamellar bonding of ceramic coatings by particle erosiontest, Thermal Spray, Current Status and Future Trends, ed.A. Ohmori, Japan High Temperature Society, 967–972.

Li C.-J., Ohmori A., and Harada Y. 1996a. Effect of powderstructure on the structure of thermally sprayed WC-Cocoatings. J. Mater. Sci. 31:785–794.

Li C.-J., Ohmori A., and Harada Y. 1996b. Formation of anamorphous phase in thermally sprayed WC-Co. J. Therm.Spray Technol. 5:69–73.

Li C.-J., Ohmori A., and McPherson R. 1992. Amicrostructural model for the elastic behavior of plasmasprayed ceramic coatings. In Proceedings of Austceram’92,(ed.) M.J. Bannister, Melbourne, August, pp. 816–821.

Li C.-J., Ohmori A., and McPherson R. 1997. Therelationship between microstructure and Young’s modulus ofthermally sprayed ceramic coatings. J. Mater. Sci.32:997–1004.

Li C.-J., Ohmori A., and Tani K. 1999a. Effect of WCparticle size on the abrasive wear of thermally sprayedWC-Co coatings. Mater. Manuf. Process. 14:175–184.

Li C.-J., and Sun B. 2004a. Effects of spray parameters onthe microstructure and property of Al 2 O 3 coatingssprayed by a low power plasma torch with a novel hollowcathode. Thin Solid Films 450:282–289. For cat logingpurposes nly

Li C.-J., and Sun B. 2004b. Microstructure and propertiesof molybdenum coatings deposited by microplasma sprayingtorch with a novel hollow cathode. In TagungsbandConference Proceedings (2004 Int. Thermal Spray Conf.)(ed.) E. Lugscheider, German Welding Research Institute,Dusseldolf, pp. 625–631.

Li C.-J., and Wang W.-Z. 2004. Quantitativecharacterization of lamellar microstructure ofplasma-sprayed alumina coating by visualization of voidsdistribution. Mater. Sci. Eng. A 386:10–19.

Li C.-J., Wang Y.-Y., and Li H. 2004a. Effect ofnano-crystallization of HVOF-sprayed NiCrBSi alloy onproperties of coatings. J. Vaccum Sci. Technol. A22:2000–2004.

Li C.-J., Wang W.-Z., and He Y. 2004c. Dependency offracture toughness of plasma-spray Al 2 O 3 coatings onthe microstructure, J. Therm. Spray Technol. 13:425–431.

Li C.-J., Wang Y.-Y., Wu T., and Ji G.-C. 2001. Effect oftypes of ceramic materials in aggregated powder on theadhesive strength of HVOF cermet coating. Surf. Coat.Technol. 145:113–120.

Li C.-J., Wang Y.-Y., Yang G.-J., Ohmori A., and Khor K. A.2004b. Effect of solid carbide particle size on depositionbehavior, microstructure and wear performance of HVOFcermet coatings. Mater. Sci. Technol. 20:1087–1096.

Li C.-J., Yang H., and Li H. 1999b. Effect of gaseousconditions on HVOF �ame and property of WC-Co coatings.Mater. Manuf. Process.14:383–395.

Li C.-J., Yang G.-J., and Ohmori A. 2003a. Potentialstrengthening of erosion performance of plasma-sprayed Al2 O 3 coating by adhesives impregnation. J. Mater. Sci.Lett. 22:1499–1501.

Li C.-J., Yang G.-J., and Ohmori A. 2003b. Improvement ofthe properties of thermally sprayed ceramic coating by thein�ltration of the adhesives. In Thermal Spray: Advancingthe Science and Applying the Technology (eds.) C. Moreauand B. Marple, ASM international, Materials Park, OH, USA,pp. 1311–1316.

Li C.-J., Yang G.-J., and Ohmori A. 2006b. Relationshipbetween particle erosion and lamellar microstructure forplasma sprayed alumina coatings. Wear 260:1166–1172.

Li C.-J., Yang G.-J., Gao P.-H., Ma J., Wang Y.-Y., and LiC.-X. 2007. Characterization of nanostructured WC-Codeposited by cold spraying. J. Therm. Spray Technol.16:1011–1012.

Liao H.-L., Zhu Y.-L., Bolot R., Coddet C., and Ma S.-N.2005. Size distribution of particles from individual wires and the effects of nozzle geometry in twin arcspraying. Surf. Coat. Technol. 200:2123–2130.

Liu X.Y., Paul K.C., and Ding C.X. 2004. Surfacemodi�cation of titanium, titanium alloys, and relatedmaterials for biomedical applications. Mater. Sci. Eng. R47:49–121.

Lima R.S., and Marple B.R. 2007. Thermal spray coatingsengineered from nanostructured ceramic agglomerated powdersfor structural, thermal barrier and biomedicalapplications: A review. J. Therm. Spray Technol. 16:40–63.

Loeb A.L. 1954. Thermal conductivity: VIII, A theory ofthermal conductivity of porous materials. J. Amer. Ceram.Soc. 37(2):96–99.

Lovelock H.L. 1998. Powder/processing/structurerelationships in WC-Co thermal spray coatings: A review ofthe published literature. J. Therm. Spray Technol.7:357–373.

Lugscheider E., Herbst C., and Zhao L.D. 1998. Parameterstudies on high-velocity oxy-fuel spraying of MCrAlYcoatings. Surf. Coat. Technol. 108–109:16–23.

Maev R. Gr., and Leshchynsky V. 2008. Introduction to LowPressure Gas Dynamic Spray, Physics and Technology,Wiley-VCH, Weinheim.

Matsumoto M., Wada K., Yamaguchi N., Kato T., and MatsubaraH. 2008. Effects of substrate rotation speed duringdeposition on the thermal cycle life of thermal barriercoatings fabricated by electron beam physical vapordeposition. Surf. Coat. Technol. 202:3507–3512.

McPherson R. 1984. A model for the thermal conductivity ofplasma-sprayed ceramic coatings. Thin Solid Films112:89–95.

McPherson R. 1988. The microstructure of thermally sprayedcoatings. Proceedings of International Symposium onAdvanced Thermal Spraying Technology and Allied Coatings(ATTAC’88), Japan High Temperature Society, Osaka, 1988,pp. 25–30.

McPherson R., and Cheang P. 1991. Elastic anisotropy of APSalumina coatings and its relationship to microstructure. InHigh Performance of Ceramic Films and Coatings (ed.) P.Vincenzini, Elsevier, Montecatini Terme, Italy, pp.277–290.

McPherson R., and Shafer B.V. 1982. Interlamellar contactwithin plasma-sprayed coatings. Thin Solid Films97:201–204.

Madejski J. 1976. Solidi�cation of droplets on a coldsurface. Int. J. Heat Mass Transfer 19:1009–1013.

Maruo H., Hirata Y., Kato J., and Matsumoto Y.1988.Development of new plasma spraying torch with threeelectrodes. Proceedings of International Symposium onAdvanced Thermal Spraying Technology and Allied Coatings(ATTAC’88), Japan High Temperature Society, Osaka, 1988,pp. 153–160.

Marple B.R., Voyer J., and Simard S. 1999. Tungstencarbide-based coatings as alternatives to electrodepositedhard chromium. Proceedings of United Thermal SprayConference 1999, DVS-German Welding Research, Dusseldolf,pp. 122–127.

Matejicek J., Kolman B., Dubsky J., Neufuss K., Hopkins N.,and Zwick J. 2006. Alternative methods for determination ofcomposition and porosity in abradable materials. Mater.Character. 57:17–29.

Miller R.A. 1997. Thermal barrier coatings for aircraftengines—history and direction. J. Therm. Spray Technol.6:35–42.

Mills A.F. 1999. Heat Transfer. Prentice Hall, Upper SaddleRiver, New Jersey, p. 8.

Miyase J. 1998. Handbook of Thermal Spraying. Japan ThermalSpray Association, Tokyo, Japan, p. 132.

Moreauo C., Cielo P., and Lamontagne M. 1992. Flatteningand solidi�cation of thermally sprayed particles.J. Therm. Spray Technol. 1:317–324.

Nakahira H., Tani K., Miyajima K., and Harada Y. 1992.Anisotropy of thermally sprayed coatings. In Proceedingsof 13th ITSC (ed.) C.C. Berndt, ASM International,Materials Park, OH USA, pp. 1011–1018.

Neuer G., Steffens H.-D., Brandl W., Babiak Z., and BrandtR. 1991. Some aspects of properties design ofplasma-sprayed thermal barrier coatings. Powder Metall.Int. 23(2):108–111.

Nicholls J.R., Lawson K.J., Johnstone A., and Rickerby D.S.2002. Methods to reduce the thermal conductivity of EB-PVDTBCs. Surf. Coat. Technol. 151–152:383–391.

Ohmori A. and Li C.-J. 1991. Quantitative characterizationof the structure of plasma sprayed Al 2 O 3 coating byusing copper electroplating. Thin Solid Films 201:241–252.

Ohmori A., and Li C.-J. 1993. The structure of thermallysprayed ceramic coatings and its dominant effect on thecoating properties. In Plasma Spraying, Theory andApplications (ed.) S. Surayanarayanan, World Scienti�cPublishing Co., Singapore, 179–200.

Ohmori A., Li C.-J., and Arata Y. 1990a. In�uence of plasmaspray conditions on the structure of Al 2 O 3 coatings.Trans. Jpn. Weld. Res. Inst. 19:259–270.

Ohmori A., Li C.-J., and Arata Y. 1992. Structure ofplasma-sprayed alumina coatings revealed by using copperelectroplating. In Thermal Spray Coating, Properties,Processes and Applications (ed.) T.F. Bernecki, ASMInternational, Materials Park, OH, USA, pp. 105–113.

Ohmori A., Li C.-J., Arata Y., Inoue K., and Iwamoto N.1990b.The dependency of the connected porosity inplasma-sprayed Al 2 O 3 coatings on microstructure. J.Jpn. High Temp. Soc.16 (Supplement):332–340 (in Japanese).

Onaran K. 2009. Malzeme Bilimi. Bilim Teknik Yayınevi,Istanbul, Turkey, pp. 171–172.

Parco M., Zhao L., Zwick J., Bobzin K., and Lugscheider E.2007. Investigation of particle �attening behaviour andbonding mechanisms of APS sprayed coatings on magnesiumalloys. Surf. Coat. Technol. 201:6290–6296.

Patankar S. V. 1997. Numerical Heat Transfer and FluidFlow. McGraw-Hill Book Company, New York, USA, pp. 3–52.

Patel, R.R., Keshri, A.K., Dulikravich, G.S., and Agarwal,A. 2010. An experimental and computational methodology fornear net shape fabrication of thin walled ceramicstructures by plasma spray forming. J. Mater. Process.Technol. 210:1260–1269.

Pawlowski L. 1995. The Science and Engineering of ThermalSpray Coatings. John Wiley & Sons Ltd., Chichester,England.

Portinha A., Teixeira V., Carneiro J., Martins J., CostaM.F., Vassen R., and Stoever D. 2005 Characterization ofthermal barrier coatings with a gradient in porosity. Surf.Coat. Technol. 195:245–251.

Raghavan S. 1998. The effect of grain size, porosity andyttria content on the thermal conductivity of

nanocrystalline zirconia. Scripta Mater. 39:1119–1125.

Rangaswamy S., and Herman H. 1986. Metallurgicalcharacterization of plasma sprayed WC-Co. In Advanced inThermal Spray, Welding Institute of Canada, Pergamon Press,New York, USA, pp. 101–110.

Rätzer-Scheibe H.-J., Schulz U., and Krell T., The effectof coating thickness on the thermal conductivity of EB-PVDPYSZ thermal barrier coatings, Surface &Coating Technology200:5636–5639.

Reed R.C. 2006. The Superalloys Fundementals &Applications. Cambridge University Press, Cambridge, UK,p. 286.

Refke A., Hawley D., Doesburg J., and Schmid R.K. 2005.LPPS thin �lm technology for the application of TBCsystems. In Proc. Int. Thermal Spray Conf. 2005 (ed.) E.Lugscheider and D. von Hofe, Basel, DVSGerman WeldingInstitute, Dusseldorf, Germany, pp. 438–443.

Rice R.W. 1977. Microstructure dependence of mechanicalbehavior of ceramics. In Properties and Microstructure(Treatise on Materials Science and Technology), Vol. 11,(ed.) R.K. MacCrone, Academic Press, New York, USA, pp.199–381.

Ruhl R.C. 1967. Cooling rates in splat cooling. Mater. Sci.Eng. 1:313–320. or cataloging purposes only

Saruhan B., Ryukhtin V., and Kelm K. 2011. Correlation ofthermal conductivity changes with anisotropic nano-poresof EB-PVD deposited FYSZ-coatings. Surf. Coat. Technol.205:5369–5378.

Simpson L.A. 1974. In Fracture Mechanics (ed.) R. C. Bradtet al., Plenum Press, NY, p. 567.

Smith M.F., McGuf� D.T., Hen�ing J.A., and Lenling W.B.1992. A comparison of techniques for the metallographicpreparation of thermal sprayed samples. In Thermal SprayCoating, Properties, Processes and Applications (ed.) T.F.Bernecki, ASM International, Materials Park, OH, USA, pp.97–104.

Sobhan C.B., and Peterson G.P. 2008. Microscale andNanoscale Heat Transfer, Fundamentals and EngineeringApplications. CRC Press, Boca Raton, FL.

Sorai M. 2004 Comprehensive Handbook of Calorimetry andThermal Analysis. John Wiley & Sons, England, pp. 167–173.

Speyer R.F. 1994. Thermal Analysis on Materials. MarcelDekker Inc, New York, USA, p. 245.

Stewart D.A., Shipway P.H., and McCartney D.G. 1999.Abrasive wear behaviour of conventional and nanocompositeHVOF-sprayed WC–Co coatings. Wear 225:789–798.

Steffens H.D., and Fischer U. 1989. Correlation betweenmicrostructure and physical properties of plasma sprayedzirconia coatings. Proceedings of 2nd National ThermalSpray Conference, ASM International, pp. 167–173.

Suryanarayanan R. 1993. Plasma Spraying: Theory andApplications. World Scienti�c Publishing Co., Singapore,1993.

Subrahmanyam J., Srivastava M.P., and Sivakumar R. 1996.Characterization of plasam-sprayed WC-Co coating. Mater.Sci. Eng. 84:209–214.

Tanner B. 2008. Thermal spraying. Sulzer Metco Product &Services.

Taylor R.E., Wang X., and Xu X. 1999. Thermophysicalproperties of thermal barrier coatings. Surf. Coat.Technol. 120–112:89–95.

Tucker Jr. R.C. 1974. Structure property relationships indeposits produced by plasma spray and detonation guntechniques. J. Vac. Sci. Technol. 11:725–734.

Tucker Jr. R.C., and Nitta H. 1988. Detonation guncoatings—Coating characteristics and applications.Proceedings of International Symposium on Advanced ThermalSpraying Technology and Allied Coatings (ATTAC’88), JapanHigh Temperature Society, Osaka, 1988, pp. 85–93.

Turunen E., Varis T., Hannula S.-P., Vaidya A., KulkarniA., Gutleber J., Sampath S., and Herman H. 2006. On therole of particle state and deposition procedure onmechanical, tribological and dielectric response of highvelocity oxy-fuel sprayed alumina coatings. Mater. Sci.Eng. A, A415: 1–11.

Vassen R., Jarligo M.O., Ophelia M., Steinke T., Mack D.E.,and Stover D. 2010. Overview on advanced thermal barriercoatings. Surf. Coat. Technol. 205:938–942.

Vassen R., Pracht G., and Stoever D. 2002. New thermalbarrier coating systems with a graded ceramic coatings. InTagungsband Conference Proceedings (ed.) E. Lugscheider,German Welding Society, pp. 202–206.

Vardelle M., and Besson J.L. 1981. γ-Alumina obtained byarc plasma spraying: A study of the optimization ofspraying conditions. Ceram. Inter. 7:48–54.

Vardelle M., Vardelle A., and Fauchais P. 1983. Study oftrajectories and temperatures of powders in a D.C. plasmajet—Correlation with alumina sprayed coatings. Proc. 10thInt. Thermal Spraying Conf., Essen, May 1983, GermanWelding Society, pp. 88–92.

Vardelle M., Vardelle A., Leger A.C., Fauchais P., andGobin D. 1995. In�uence of particle parameters at impacton splat formation and solidi�cation in plasma sprayingprocesses. J. Therm. Spray Technol. 4:50–58.

Vardelle A., Vadelle M., McPherson R., and Fauchais P.1980. Study of the in�uence of particle temperature andvelocity distribution within a plasma jet coatingformation. Proc. 9th Int. Thermal Spraying Conf, Hague,1980, Nederland Inst. Voor Lastetechnik, pp. 155–161.

Vinayo M.E., Kassabji F., Guyonnet J., and Fauchais P.1985. Plasma sprayed WC-Co coatings—In�uence of sprayconditions (atmospheric and low-pressure plasma spraying)on the crystal-structure, porosity, and hardness. J. Vac.Sci. Technol. A 3:2483–2489.

von Niessen K., Gindra M., and Refke A. 2010. Vapor phasedeposition using plasma sprayed-PVD TM . J. Therm. SprayTechnol. 19:502–509.

Wagner N., Gnadig K., Kreye H., and Kronewetter H. 1984.Particle velocity in hypersonic �ame spraying of WC-Co.Furf. Technol. 22:61–71.

Walser B. 2004. The importance of thermal spray for currentand future applications in key industries. Spraytime10(4):1–7.

Wang Y.-Y., and Li C.-J. 2005. Effect of substrate surfaceconditions on the bonding mechanism of HVOF coatings. ThinSolid Films 485:141–147.

Wang Y.-Y., Li C.-J., Ma J., and Yang G.-J. 2004. Effect of

�ame conditions on the abrasive wear performance of HVOFsprayed nanostructured WC-12Co coatings. Trans. NonferrousMet. Soc. China 14(special 2):72–76.

Wang W.-Z., Li C.-J., and Sonoya K. 2005. Study of lamellarmicrostructure of plasma-sprayed ZrO 2 -8wt.%Y 2 O 3coatings. In Tagungsband Conference Proceedings (2005 Int.Thermal Spray Conf.) (ed.) E. Lugscheider, May 2–5,Bussel, Switzerland, German Welding Research Institute,Germany, pp. 1506–1510.

Wang Y.Y., Li C.J., Yang G.J., and Li Q.L. 2004b. Effect of�bre volume fractions on the mechanical performance ofB/Al composite fabricated by plasma spraying. AeronauticalManuf. Technol. Special/2004:285–288 (in Chinese).

Wayne S.F., Baldoni J.G., and Buljan S.-T. 1990. Abrasionand erosion of WC-Co with controlled microstructure.Tribol. Trans. 33:611–617.

Whittmore O.J. 1981. Mercury porosimetry of ceramics.Powder Technol. 29:167–175.

Wu J., Guo H., Zhou L., Wang L., and Gong S. 2010.Microstructure and thermal properties of plasma sprayedthermal barrier coatings from nanostructured YSZ. J. Therm.Spray Technol. 19:1186–1194.

Xie X.Y., Guo H., and Gong S. 2010. Mechanical propertiesof LaTi 2 Al 9 O 19 and thermal cycling behaviors ofplasma-sprayed LaTi2Al9O19/YSZ thermal barrier coatings. J.Therm. Spray Technol. 19:1179–1185.

Xie X., Guo H., Gong S., and Xu H. 2011.Lanthanum–titanium–aluminum oxide: A novel thermal barriercoating material for applications at 1300°C. J. Eur. Ceram.Soc. 31:1677–1683.

Xing Y.-Z., Li C.-J., Li C.-X., and Yang G.-J. 2008a.In�uence of through-lamella grain growth on ionicconductivity of plasma-sprayed yttria stabilized zirconiaas an electrolyte in solid oxide fuel cells. J. PowerSources 176:31–38.

Xing Y.-Z., Li C.-J., Li C.-X., and Yang G.-J. 2008b.In�uence of substrate temperature on microcracks formationin plasma-sprayed yttria-stabilized zirconia splats. KeyEng. Mater. 373–374:69–72.

Xing Y.-Z., Li C.-J., Zhang Q., Li C.-X., and Yang G.-J.

2008c. In�uence of microstructure on the ionic conductivityof plasma-sprayed Yttria-stabilized Zirconia deposits. J.Amer. Ceram. Soc. 91:3931–3936.

Xu B., Zhu Z., Ma S., Zhang W., and Liu W. 2004. Slidingwear behavior of Fe-Al and Fe-Al/WC coatings prepared byhigh velocity arc spraying. Wear 257:1089–1095.

Yang G.-J., Li C.-J., Han F., Li W.-Y., and Ohmori A. 2008.Low temperature deposition and characterization of TiO 2photocatalytic �lm through cold spray. Appl. Surf. Sci.54:3979–3982.

Yang Q., Senda T., and Ohmori A. 2003. Effect of carbidegrain size on microstructure and sliding wear behavior ofHVOF-sprayed WC-12% Co coatings. Wear 254:23–34.

Yi M.Z., He J.W., Huang B.Y., and Zhou H. 1999. Frictionand wear behaviour and abradability of abradable sealcoating. Wear 231:47–53.

Yoshida T. 2006. Toward a new era of plasma sprayprocessing. Pure Appl. Chem. 78:1093–1107.

Yoshiya M., Wada K., Jang B.K., and Matsubara H. 2004.Computer simulation of nano-pore formation in EB-PVDthermal barrier coatings. Surf. Coat. Technol. 187:399–407.

Yüncü H., and Kakaç S. 1999. Temel Isı Transferi. BilimYayıncılık, Ankara, Turkey, p. 3.

Zhu Y.C., Yukimura K., Ding C.X., and Zhang P.Y. 2001.Tribological properties of nanostructured and conventionalWC-Co coatings deposited by plasma spraying. Thin SolidFilms 388:277–282.

Zimmermann S., and Kreye H. 1996. Chromium carbide coatingsproduced with various HVOF spray systems. Thermal Spray:Practical Solutions For Engineering Problems, 9th NationalThermal Spray Conference & Exposition, 1996 Cincinnati,ASM International, Ohio, pp. 147–152. For cataloging prposes only

7 Chapter 7 - Nanostructured SolidLubricant Coatings for AerospaceApplications

1. Bhushan, B. 1999. Principles and Applications ofTribology. New York: John Wiley & Sons, Inc.

2. Ludema, K. C. 2001. Friction. In Modern TribologyHandbook, Volume 1, ed. B. Bhushan, pp. 205–233. BocaRaton: CRC Press LLC.

3. Bowden, F. P. and Tabor, D. 1950. The Friction andLubrication of Solids. London: Oxford University Press.

4. Mo, Y., Turner, K. T., and Szlufarska, I. 2009. Frictionlaws at the nanoscale. Nature 457: 1116–1119.

5. Moore, D. F. 1975. Principles and Applications ofTribology. Oxford: Pergamon Press.

6. Kato, K. and Adachi, K. 2001. Wear mechanisms. In ModernTribology Handbook, Volume 1, ed. B. Bhushan, pp. 273–300.Boca Raton: CRC Press LLC.

7. Jones Jr. W. R. and Jansen, M. J. 2000. Space tribology.NASA/TM-2000-209924.

8. Roberts, E. W. and Todd, M. J. 1990. Space and vacuumtribology. Wear 136: 157–167.

9. Clauss, F. J. 1972. Solid Lubricants and SelfLubricating Solids. New York: Academic Press.

10. Campbell, M. E. 1972. Solid lubricants—A survey. NASASP-5059(01).

11. Sliney, H. E. 1991. Solid lubricants. NASA TM 103803.

12. McMurtrey, E. L. 1985. Lubrication handbook for thespace industry. NASA TM-86556.

13. Miyoshi, K. 2007. Solid lubricants and coatings forextreme environments: State-of-the-art survey.NASA/TM-2007-214668.

14. Spalvins, T. 1970. Friction characteristics ofsputtered solid �lm lubricants. NASA TM X-52819.

15. Miyoshi, K., Iwaki, M., Gotoh, K., Obara, S., andImagawa, K. 1999. Friction and wear properties of selected

solid lubricant �lms. NASA TM-1999-209088.

16. Sliney, H. E. 1973. High temperature solid lubricant�lms—When and where to use them. NASA TM X-68201.

17. Donnet, C. and Erdemir, A. 2004. Historicaldevelopments and new trends in tribological and solidlubricant coatings. Surf. Coat. Technol. 180–181: 76–84.

18. Braithwaite, E. R. 1964. Solid Lubricants and Surfaces.Oxford: Pergamon Press.

19. Erdemir, A. 2001. Solid lubricants and self-lubricating�lms. In Modern Tribology Handbook, Vol. 2, ed. B.Bhushan, pp. 787–825. Boca Raton: CRC Press LLC.

20. Hilton, M. R. and Fleischauer, P. D. 1992. Applicationsof solid lubricant �lms in spacecraft. Surf. Coat.Technol. 54–55: 435–441.

21. Stachowiak, G. W. and Batchelor, A. W. 2005.Engineering Tribology. Oxford: ElsevierButterworth-Heinemann.

22. Chopra K. L. 1969. Thin Film Phenomena. New York:McGraw Hill.

23. Savage, R. H. 1948. Graphite lubrication. J. Appl.Phys. 19: 1–10.

24. Deacon. R. F. and Goodman, J. F. 1958. Lubrication bylamellar solids. Proc. R. Soc. Lond. A 243: 464–482.

25. Yen, B. K., Schwickert, B. E., and Toney, M. F. 2004.Origin of low-friction behavior in graphite investigated bysurface x-ray diffraction. Appl. Phys. Lett. 84: 4702–4704.

26. Spalvins, T. 1987. A review of recent advances in solid�lm lubrication. J. Vac. Sci. Technol. A 5: 212–219.

27. Donnet, C., Martin, J. M., Mogne, T. L., and Belin, M.1996. Super-low friction of MoS 2 coatings in variousenvironments. Tribol. Int. 29: 123–128.

28. Fleischauer, P. D. and Lince, J. R. 1999. A comparisonof oxidation and oxygen substitution in MoS 2 solid �lmlubricants. Tribol. Int. 32: 627–636.

29. Erdemir, A. 1994. Crystal chemistry and solidlubricating properties of themonochalcogenides gallium

selenide and tin selenide. Tribol. Trans. 37: 471–478.

30. Martin, J. M., Mogne, T. L., Chassagnette, C., andGardos, M. N. 1992. Friction of hexagonal boron nitride invarious environments. Tribol. Trans. 35: 462–472.

31. Li, Y., Ruoff, R. S., and Chang, R. P. H. 2003. Boricacid nanotubes, nanotips, nanorods, microtubes, andmicrotips. Chem. Mater. 15: 3276–3285.

32. Erdemir, A., Bindal, C., and Fenske, G. R. 1996.Formation of ultralow friction surface �lms on boroncarbide. Appl. Phys. Lett. 68: 1637–1639.

33. Robertson, W. M. 1981. Sodium chromite as afriction-reducing material. Wear 68: 97–107.

34. Fusaro, R. L. 1979. Mechanisms of graphite �uoride [(CFx ) n ] lubrication. Wear 53: 303–323.

35. Holmberg, K. and Matthews, A. 2009. Coatings Tribology.Oxford: Elsevier.

36. Sherbiney, M. A. and Halling, J. 1977. Friction andwear of ion-plated soft metallic �lms. Wear 45: 211–220.

37. Rabinowicz, E. 1967. Variation of friction and wear ofsolid lubricant �lms with �lm thickness. Tribol. Trans.10: 1–9.

38. Ravindran, K. A., Ramasamy, P., and Laddha, G. S. 1980.Frictional behavior of thin �lms of soft metals. ThinSolid Films 66: 249–254.

39. Arnell, R. D. and Soliman, F. A. 1978. The effects ofspeed, �lm thickness and substrate surface roughness on thefriction and wear of soft metal �lms in ultrahigh vacuum.Thin Solid Films 53: 333–341.

40. Spalvins, T. and Buzek, B. 1981. Frictional andmorphological characteristics of ion-plated soft metallic�lms. Thin Solid Films 84: 267–272.

41. Miyoshi, K. 2001. Solid Lubrication Fundamentals andApplications. New York: Marcel Dekker, Inc.

42. Erdemir, A. and Donnet, C. 2001. Tribology of diamond,diamond-like carbon and related �lms. In Modern TribologyHandbook, Vol. 2, ed. B. Bhushan, pp. 871–908. Boca Raton:CRC Press.

43. Robertson, J. 2002. Diamond-like amorphous carbon.Mater. Sci. Eng. R 37:129–281.

44. Fontaine, J., Donnet C., and Erdemir, A. 2008.Fundamentals of the tribology of DLC coatings. InTribology of Diamond-Like Carbon Films: Fundamentals andApplications, ed. C. Donnet and A. Eredmir, pp. 139–154.New York: Springer.

45. Vanhulsel, A., Velasco, F., Jacobs, R. et al. 2007. DLCsolid lubricant coatings on ball bearings for spaceapplications. Tribol. Int. 40: 1186–1194.

46. Grill, A. 1993. Review of the tribology of diamond-likecarbon. Wear 168: 143–153.

47. Erdemir, A. and Eryilmaz, O. L. 2007. Superlubricity indiamondlike carbon �lms. In Superlubricity, ed. A. Erdemirand J.-M. Martin, pp. 253–271. Amsterdam: Elsevier.

48. Liu, Y., Erdemir, A., and Meletis, E. I. 1997. In�uenceof environmental parameters on the frictional behavior ofDLC coatings. Surf. Coat. Technol. 94–95: 463–468.

49. Erdemir, A. and Donnet, C. 2006. Tribology ofdiamond-like carbon �lms: Recent progress and futureprospects. J. Phys. D: Appl. Phys. 39: R311–R327.

50. Ronkainen, H. and Holmberg, K. 2008. Environmental andthermal effects on the tribological performance of DLCcoatings. In Tribology of Diamond Like Carbon Films:Fundamentals and Applications, ed. C. Donnet and A.Eredmir, pp. 155–200. New York: Springer. For cat l gingpurposes only

51. Erdemir, A. 2004. Design criteria for superlubricity incarbon �lms and related microstructures. Tribol. Int. 37:577–583.

52. Erdemir, A., Eryilmaz, O. L., and Fenske, G. 2000.Synthesis of diamondlike carbon �lms with superlowfriction and wear properties. J. Vac. Sci. Technol. A 18:1987–1992.

53. Erdemir, A. 2000. A crystal-chemical approach tolubrication by solid oxides. Tribol. Lett. 8: 97–102.

54. Goto, M., Kasahara, A., Oishi, T., Konishi, Y., andTosa, M. 2003. Lubricative coatings of copper oxide for

aerospace applications. J. Appl. Phys. 94: 2110–2114.

55. Zabinski, J. S., Sanders, J. H., Nainaparambil, J., andPrasad, S. V. 2000. Lubrication using a microstructurallyengineered oxide: Performance and mechanisms. Tribol. Lett.8: 103–116.

56. Zabinski, J. S., Day, A. E., Donley, M. S., Dellacorte,C., and McDevitt, N. T. 1994. Synthesis andcharacterization of a high-temperature oxide lubricant. J.Mater. Sci. 29: 5875–5879.

57. Allam, I. M. 1991. Solid lubricants for applications atelevated temperatures. J. Mater. Sci. 26: 3977–3984.

58. Buckley, D. H. 1981. Surface Effects in Adhesion,Friction, Wear, and Lubrication. Amsterdam: Elsevier.

59. Ma, K. J., Chao, C. L., Liu, D. S., Chen, Y. T., andShieh, M. B. 2002. Friction and wear behaviour of TiN/Au,TiN/MoS 2 and TiN/TiCN/a-C:H coatings. J. Mater. Proc.Technol. 127: 182–186.

60. Shtansky, D. V., Lobova, T. A., Yu, V. et al. 2004.Structure and tribological properties of WSe x , WSe x /TiN, WSe x /TiCN and WSe x /TiSiN coatings. Surf. Coat.Technol. 183: 328–336.

61. Cselle, T. and Barimani, A. 1995. Today’s applicationsand future developments of coatings for drills androtating cutting tools. Surf. Coat. Technol. 76–77:712–718.

62. Musil, J. 2006. Physical and mechanical properties ofhard nanocomposite �lms prepared by reactive magnetronsputtering. In Nanostructured Coatings, ed. A. Cavaleiro,and J. T. M. De Hosson, pp. 407–463. New York: Springer.

63. Barshilia, H. C., Deepthi, B., and Rajam, K. S. 2010.Transition metal nitride-based nanolayered multilayercoatings and nanocomposite coatings as novel superhardmaterials. In Handbook of Nanostructured Thin Films andCoatings: Mechanical Properties, ed. S. Zhang, pp. 427–480.Boca Raton: CRC Press.

64. Tjong, S. C. and Chen, H. 2004. Nanocrystallinematerials and coatings. Mater. Sci. Eng. R 45: 1–88.

65. Erdemir, A. and Voevodin, A. A. 2010. Nanocompositecoatings for severe applications. In Handbook of

Deposition Technologies for Films and Coatings: Science,Applications and Technology, ed. P. M. Martin, pp.679–715. Oxford: Elsevier.

66. Jehn, H. A. 2000. Multicomponent and multiphase hardcoatings for tribological applications. Surf. Coat.Technol. 131: 433–440.

67. Savan, A., P�üger, E., Goller, R., and Gissler, W.2000. Use of nanoscaled multilayer and compound �lms torealize a soft lubrication phase within a hard,wear-resistant matrix. Surf. Coat. Technol. 126:159–165.

68. Holleck, H. 1999. Design of nanostructured thin �lmsfor tribological applications. In Surface Engineering:Science and Technology I, ed. A. Kumar, Y.-W Chung, J. J.Moore, and J. E. Smugeresky, pp. 207–218. Pennsylvania:The Minerals, Metals and Materials Society.

69. Yashar, P. C. and Sproul, W. D. 1999. Nanometer scalemultilayered hard coatings. Vacuum 55: 179–190.

70. Koehler, J. S. 1970. Attempt to design a strong solid.Phys. Rev. B 2: 547–551.

71. Voevodin, A. A., Zabinski, J. S., and Muratore, C.2005. Recent advances in hard, tough, and low frictionnanocomposite coatings. Tsinghua Sci. Technol. 10: 665–679.

72. Barshilia, H. C., Jain, A., and Rajam, K. S. 2004.Structure, hardness and thermal stability of nanolayeredTiN/CrN multilayer coatings. Vacuum 72: 241–248.

73. Barshilia, H. C., Deepthi, B., Selvakumar, N., Jain A.,and Rajam, K. S. 2007. Nanolayered multilayer coatings ofCrN/CrAlN prepared by reactive DC magnetron sputtering.Appl. Surf. Sci. 253: 5076–5083.

74. Vepřek, S., Reiprich, S., and Shizhi, L. 1995.Superhard nanocrystalline composite materials: The TiN/ Si3 N 4 system. Appl. Phys. Lett. 66: 2640–2642.

75. Musil, J. and Vlček. J. 2001. Magnetron sputtering ofhard nanocomposite coatings and their properties. Surf.Coat. Technol. 142–144: 557–566.

76. Barshilia, H. C., Prakash, M. S., Rao, D. V. S., andRajam, K. S. 2005. Superhard nanocomposite coatings ofTiN/a-C prepared by reactive DC magnetron sputtering. Surf.Coat. Technol. 195: 147–153.

77. Barshilia, H. C., Deepthi, B., Arun Prabhu, A. S., andRajam, K. S. 2006. Superhard nanocomposite coatings ofTiN/Si 3 N 4 prepared by reactive direct currentunbalanced magnetron sputtering. Surf. Coat. Technol. 201:329–337.

78. Barshilia, H. C., Deepthi, B., and Rajam, K. S. 2007.Deposition and characterization of CrN/Si 3 N 4 andCrAlN/Si 3 N 4 nanocomposite coatings prepared usingreactive DC unbalanced magnetron sputtering. Surf. Coat.Technol. 201: 9468–9475.

79. Lehoczky, S. L. 1978. Retardation of dislocationgeneration and motion in thin layered metal laminates.Phys. Rev. Lett. 41: 1814–1818.

80. Barnett, S. A. and Shinn, M. 1994. Plastic and elasticproperties of compositionally modulated thin �lms. Annu.Rev. Mater. Sci. 24: 481–511.

81. Chu, X. and Barnett, S. A. 1995. Model of superlatticeyield stress and hardness enhancements. J. Appl. Phys. 77:4403–4411.

82. Barnett, S. A. and Madan, A. 1998. Superhardsuperlattices. Phys. World. January issue: 45–50.

83. Barnett, S. A. 1993. Deposition and mechanicalproperties of superlattice thin �lms. In Physics of ThinFilms, Vol. 17, ed. M. H. Francombe and J. L. Vossen, 1–77.Boston: Academic Press, Inc.

84. Hilton, M. R., Bauer, R., Didziulis, S. V., Dugger, M.T., Keem, J. M., and Scholhamer, J. 1992. Structural andtribological studies of MoS 2 solid lubricant �lms havingtailored metal-multilayer nanostructures. Surf. Coat.Technol. 53: 13–23.

85. Simmonds, M. C., Savan, A., Swygenhoven, H. V., P�üger,E., and Mikhailov, S. 1998. Structural, morphological,chemical and tribological investigations of sputterdeposited MoS x /metal multilayer coatings. Surf. Coat.Technol. 108–109: 340–344.

86. Mikhailov, S., Savan, A., P�üger, E. et al. 1998.Morphology and tribological properties of metal (oxide)MoS2 nanostructured multilayer coatings. Surf. Coat. Technol.105: 175–183.

87. Jayaram, G., Marks, L. D., and Hilton, M. R. 1995.Nanostructure of Au-20%Pd layers in MoS 2 multilayer solidlubricant �lms. Surf. Coat. Technol. 76–77: 393–399.

88. Watanabe, S., Noshiro, J., and Miyake, S. 2004.Tribological characteristics of WS 2 /MoS 2 solidlubricating multilayer �lms. Surf. Coat. Technol. 183:347–351.

89. Zhang, S., Sun, D., Fu, Y., and Du, H. 2003. Recentadvances of superhard nanocomposite coatings: A review.Surf. Coat. Technol. 167: 113–119.

90. Hu, J. J., Bultman, J. E., and Zabinski, J. S. 2006.Microstructure and lubrication mechanism of multilayeredMoS 2 /Sb 2 O 3 thin �lms. Tribol. Lett. 21:169–174.

91. Shtansky, D. V., Sheveyko, A. N., Sorokin, D. I., Lev,L. C., Mavrin, B. N., and Kiryukhantsev-Korneev, P. V.2008. Structure and properties of multi-component andmultilayer TiCrBN/WSe x coatings deposited by sputteringof TiCrB and WSe 2 targets. Surf. Coat. Technol.202:5953–5961.

92. Voumard, P., Savan, A., and P�üger, E. 2001. Advancesin solid lubrication with MoS 2 multilayered coatings.Lubr. Sci. 13:135–145.

93. Martínez, E., Romero, J., Lousa, A., and Esteve, J.2003. Wear behavior of nanometric CrN/Cr multilayers. Surf.Coat. Technol. 163–164: 571–577.

94. Yamamoto, K., Ito, H., and Kujime, S. 2007.Nano-multilayered CrN/BCN coating for anti-wear and lowfriction applications. Surf. Coat. Technol. 201:5244–5248.

95. Miyake, S., Hashizume, T., Kurosaka, W., Sakurai, M.,and Wang, M. 2007. Deposition and tribology of carbon andboron nitride nanoperiod multilayer solid lubricating �lms.Surf. Coat. Technol. 202: 1023–1028.

96. Luo, Q., Schimpf, C., Ehiasarian, A. P., Chen, L., andHovsepian, P. E. 2007. Structure and wear mechanisms ofnanostructured TiAlCN/VCN multilayer coatings. PlasmaProcess. Polym. 4: S916–S920.

97. Yang, S.-M., Chang, Y.-Y., Lin, D.-Y., Wang, D.-Y., andWu, W. 2008. Mechanical and tribological properties ofmultilayered TiSiN/CrN coatings synthesized by a cathodicarc deposition process. Surf. Coat. Technol.

202:2176–2181.

98. Fox-Rabinovich, G. S., Yamamoto, K., Veldhius, S. C.,Kovalev, A. I., Shuster, L. S., and Ning, L. 2006.Self-adaptive wear behavior of nano-multilayered TiAlCrN/WNcoatings under severe machining conditions. Surf. Coat.Technol. 201:1852–1860.

99. Fox-Rabinovich, G. S., Yamamoto, K., Kovalev, A. I. et al. 2008. Wear behavior of adaptive nano- multilayeredTiAlCrN/NbN coatings under dry high performance machiningconditions. Surf. Coat. Technol. 202: 2015–2022.

100. Baraket, M., Mercs, D., Zhang, Z. G., and Coddet, C.2010. Mechanical and tribological properties of CrN/Ag andCrSiN/Ag nanoscale multilayers. Surf. Coat. Technol. 204:2386–2391.

101. Efeoglu, I. 2007. Deposition and characterization of amultilayered-composite solid lubricant coating. Rev. Adv.Mater. Sci. 15: 87–94.

102. Gulbiński, W., Suszko, T., Sienicki, W., andWarcholiński, B. 2003. Tribological properties of silver-and copper-doped transition metal oxide coatings. Wear 254(2003): 129–135.

103. Maréchal, N., Pauleau, Y., Quesnel, E., Juliet, P.,Rouzaud, A., and Zimmermann, C. 1994. Sputterdepositedlubricant thin �lms operating at elevated temperatures inair. Surf. Coat. Technol. 68–69: 416–421.

104. Voevodin, A. A., Walck, S. D., and Zabinski, J. S.1997. Architecture of multilayer nanocomposite coatingswith super-hard diamond-like carbon layers for wearprotection at high contact loads. Wear 203– 204: 516–527.For cataloging purposes only

105. Cruz, R., Rao, J., Rose, T., Lawson, K., and Nicholls,J. R. 2006. DLC-ceramic multilayers for automotiveapplications. Diamond Relat. Mater. 15: 2055–2060.

106. Yi, J. W., Kim, J.-K., Moon, M.-W., Lee, K.-R., andKim, S.-S. 2009. Tribological performance ofalternating-layered Si-DLC/DLC �lms under humid conditions.Tribol. Lett. 34: 223–228.

107. Deng, J. and Braun, M. 1995. DLC multilayer coatingsfor wear protection. Diamond Relat. Mater. 4: 936–943.

108. Dekempeneer, E., Acker, K. V., Vercammen, K. et al.2001. Abrasion resistant low friction diamond-likemultilayers. This article is not included in yourorganization’s subscription. However, you may be able toaccess this article under your organization’s agreementwith Elsevier. Surf. Coat. Technol. 142–144: 669–673.

109. Donnet, C., Fontaine, J., Mogne, T. L. et al. 1999.Diamond-like carbon-based functionally gradient coatingsfor space tribology. Surf. Coat. Technol. 120–121: 548–554.

110. Rao, J., Cruz, R., Lawson, K. J., and Nicholls, J. R.2005. Sputtered DLC-TiB 2 multilayer �lms for tribologicalapplications. Diamond Relat. Mater. 14: 1805–1809.

111. Pinyol, A., Bertran, E., Corbella, C., Polo, M. C.,and Andújar, J. L. 2002. Properties of W/a-C nanometricmultilayers produced by RF-pulsed magnetron sputtering.Diamond Relat. Mater. 11: 1000–1004.

112. Hauert, R., Knoblauch-Meyer, L., Francz, G.,Schroeder, A., and Wintermantel, E. 1999. Tailored a-C:Hcoatings by nanostructuring and alloying. Surf. Coat.Technol. 120–121: 291–296.

113. Meneve, J., Dekempeneer, E., Wegener, W., and Smeets.J. 1996. Low friction and wear resistant a-C:H/ a-Si 1−x Cx :H multilayer coatings. Surf. Coat. Technol. 86–87:617–621.

114. Kok, Y. N., Hovsepian, P. E., Luo, Q., Lewis, D. B.,Wen, J. G., and Petrov, I. 2005. In�uence of the biasvoltage on the structure and the tribological performanceof nanoscale multilayer C/Cr PVD coatings. Thin SolidFilms 475: 219–226.

115. Gilmore, R., Baker, M. A., Gibson, P. N., and Gissler,W. 1999. Comparative investigation of multilayer TiB 2 /Cand co-sputtered TiB 2 -C coatings for low-frictionapplications. Surf. Coat. Technol. 116–119: 1127–1132.

116. Rincón, C., Zambrano, G., Carvajal, A. et al. 2001.Tungsten carbide/diamond-like carbon multilayer coatingson steel for tribological applications. Surf. Coat.Technol. 148: 277–283.

117. Sobota, J., Bochnicek, Z., and Holy, V. 2003. Frictionand wear properties of C-N/MeN x nanolayer composites.Thin Solid Films 433: 155–159.

118. Hovsepian, P. E., Ehiasarian, A. P., and Ratayski, U.2009CrAlYCN/CrCN nanoscale multilayer PVD coatingsdeposited by the combined High Power Impulse MagnetronSputtering/Unbalanced Magnetron Sputtering (HIPIMS/UBM)technology. Surf. Coat. Technol. 203: 1237–1243.

119. Hu, J. J., Zabinski, J. S., Bultman, J. E., Sanders,J. H., and Voevodin, A. A. 2006. Structure characterizationof pulsed laser deposited MoS x -WSe y composite �lms oftribological interests. Tribol. Lett. 24: 127–135.

120. Savan, A., Haefke, H., Simmonds, M. C., and Constable,C. P. 2002. Microstructure, chemistry, and tribologicalperformance of MoS x /WSe y co-sputtered composited. J.Vac. Sci. Technol. A 20: 1682–1689.

121. Teer, D. G., Hampshire, J., Fox, V., andBellido-Gonzalez, V. 1997. The tribological properties ofMoS 2 / metal composite coatings deposited by closed �eldmagnetron sputtering. Surf. Coat. Technol. 94–95: 572–577.

122. Scharf, T. W., Rajendran, A., Banerjee, R., andSequeda, F. 2009. Growth, structure and friction behaviorof titanium doped tungsten disulphide (Ti-WS 2 )nanocomposite thin �lms. Thin Solid Films 517: 5666–5675.

123. Deepthi, B., Barshilia, H. C., Rajam, K. S. et al.2010. Structure, morphology and chemical composition ofsputter deposited nanostructured Cr-WS 2 solid lubricantcoatings. Surf. Coat. Technol. 205: 565–574.

124. Efeoglu, I., Baran, Ö., Yetim, F., and Altintaş, S.2008. Tribological characteristics of MoS 2 -Nb solidlubricant �lm in different tribo-test conditions. Surf.Coat. Technol. 203: 766–770.

125. Spalvins, T. 1984. Frictional and morphologicalproperties of Au-MoS 2 �lms sputtered from a compacttarget. Thin Solid Films 118: 375–384.

126. Lince J. R. 2004. Tribology of co-sputterednanocomposite Au/MoS 2 solid lubricant �lms over a widecontact stress range. Tribol. Lett. 17: 419–428.

127. Simmonds, M. C., Savan, A., P�üger, E., andSwygenhoven, H. V. 2000. Mechanical and tribologicalperformance of MoS 2 co-sputtered composites. Surf. Coat.Technol. 126: 15–24.

128. Ding, X.-Z., Zeng, X. T., He, X. Y., and Chen, Z.

2010. Tribological properties of Cr- and Ti-doped MoS 2composite coatings under different humidity atmosphere.Surf. Coat. Technol. 205: 224–231.

129. Su, Y. L. and Kao, W. H. 2003. Tribological behaviourand wear mechanism of MoS 2 -Cr coatings sliding againstvarious counterbody. Tribol. Int. 36: 11–23.

130. Nainaparampil, J. J., Phani, A. R., Krzanowski, J. E.,and Zabinski, J. S. 2004. Pulsed laser-ablated MoS 2 -Al�lms: Friction and wear in humid conditions. Surf. Coat.Technol. 187: 326–335.

131. Teer, D. G. 2001. New solid lubricant coatings. Wear251: 1068–1074.

132. Renevier, N. M., Hampshire, J., Fox, V. C., Witts, J.,Allen, T., and Teer, D. G. 2001. Advantages of usingself-lubricating, hard, wear-resistant MoS 2 -basedcoatings. Surf. Coat. Technol.142–144: 67–77.

133. Renevier, N. M., Fox, V. C., Teer, D. G., andHampshire, J. 2000. Coating characteristics andtribological properties of sputter-deposited MoS 2 -metalcomposite coatings deposited by closed �eld unbalancedmagnetron sputter ion plating. Surf. Coat. Technol. 127:24–37.

134. Teer, D. G., Bellido-Gonzalez, V., and Hampshire, J.H. 1997. United Kingdom Patent No. GB 2303380A.

135. Wang, D.-Y., Chang, C.-L., Chen, Z.-Y., and Ho, W.-Y.1999. Microstructural and tribological characterization ofMoS 2 -Ti composite solid lubricating �lms. Surf. Coat.Technol. 120–121: 629–635.

136. Savan, A., Simmonds, M. C., Huang, Y. et al. 2005.Effects of temperature on the chemistry and tribology ofco-sputtered MoS x -Ti composite thin �lms. Thin SolidFilms 489: 137–144.

137. Arslan, E., Bülbül, F., Alsaran, A., Celik, A., andEfeglu, I. 2005. The effect of deposition parameters andTi content on structural and wear properties of MoS 2 -Ticoatings. Wear 259: 814–819.

138. Stoyanov, P., Fishman, J. Z., Lince, J. R., andChromik, R. R. 2008. Micro-tribological performance of MoS2 lubricants with varying Au content. Surf. Coat. Technol.203: 761–765.

139. Lince, J. R., Kim, H. I., Adams, P. M., Dickrell, D.J., and Dugger, M. T. 2009. Nanostructural, electrical,and tribological properties of composite Au-MoS 2coatings. Thin Solid Films 517: 5516–5522.

140. Deepthi, B., Barshilia, H. C., Rajam, K. S., Konchady,M. S., Pai, D. M., and Sankar, J. 2010. Mechanical andtribological properties of sputter deposited nanostructuredCr-WS 2 solid lubricant coatings. Surf. Coat. Technol.205: 1937–1946.

141. Polcar, T., Nossa, A., Evaristo, M., and Cavaleiro, A.2007. Nanocomposite coatings of carbon-based andtransition metal dichalcogenide phases: A review. Rev. Adv.Mater. Sci. 15:118–126.

142. Pimentel, J. V., Polcar, T., and Cavaleiro, A. 2011.Structural, mechanical and tribological properties ofMo-S-C solid lubricant coating. Surf. Coat. Technol. 205:3274–3279.

143. Polcar, T., Evaristo, M., and Cavaleiro, A. 2009.Self-lubricating W-S-C nanocomposite coatings. PlasmaProcess. Polym. 6: 417–424.

144. Polcar, T., Evaristo, M., and Cavaleiro, A. 2009.Comparative study of the tribological behavior ofselflubricating W-S-C and Mo-Se-C sputtered coatings. Wear266: 388–392.

145. Polcar, T., Evaristo, M., Colaço, R., Sandu, C. S.,and Cavaleiro, A. 2008. Nanoscale triboactivity: Theresponse of Mo-Se-C coatings to sliding. Acta Mater. 56:5101–5111.

146. Panich, N., Wangyao, P., Hannongbua, S.,Sricharoenchai, P., and Sun, Y. 2007. Effect ofpolytetra�uoroethylene doping on tribological propertyimprovement of MoS 2 nano-thin �lms on Ti-substrate. Rev.Adv. Mater. Sci. 16: 88–95.

147. Gilmore, R., Baker, M. A., Gibson, P. N. et al. 1998.Low-friction TiN-MoS 2 coatings produced by dc magnetronco-deposition. Surf. Coat. Technol. 108–109: 345–351.

148. Goller, R., Torri, P., Baker, M. A., Gilmore, R., andGissler, W. 1999. The deposition of low-friction TiNMoS xhard coatings by a combined arc evaporation and magnetronsputter process. Surf. Coat. Technol. 120–121: 453–457.

149. Gangopadhyay, S., Acharya, R., Chattopadhyay, A. K.,and Paul, S. 2009. Composition and structurepropertyrelationship of low friction, wear resistant TiN-MoS xcomposite coating deposited by pulsed closed-�eldunbalanced magnetron sputtering. Surf. Coat. Technol. 203:1565–1572.

150. Cosemans P., Zhu X., Celis J. P., and Stappen, M. V.2003. Development of low friction wear-resistant coatings.Surf. Coat. Technol. 174–175:416–420.

151. Bae, Y. W., Lee, W. Y., Besmann, T. M., Yust, C. S.,and Blau, P. J. 1996. Preparation and frictioncharacteristics of self-lubricating TiN-MoS 2 compositecoatings. Mater. Sci. Eng. A 209: 372–376.

152. Deepthi, B., Barshilia, H. C., Rajam, K. S., Konchady,M. S., Pai, D. M., and Sankar, J. 2011. Structural,mechanical and tribological investigations of sputterdeposited CrN-WS 2 nanocomposite solid lubricantcoatings. Tribol. Int. 44: 1844–1851.

153. Gilmore, R., Baker, M. A., Gibson, P. N., and Gissler,W. 1998. Preparation and characterization of lowfrictionTiB 2 -based coatings by incorporation of C or MoS 2 .Surf. Coat. Technol. 105: 45–50.

154. Steinmann, M., Müller, A., and Meerkamm, H. 2004. Anew type of tribological coating for machine elementsbased on carbon, molybdenum disulphide and titaniumdiboride. Tribol. Int. 37: 879–885.

155. Efeoglu, I. 2005. Co-sputtered Mo:S:C:Ti:B basedcoating for tribological applications. Surf. Coat.Technol. 200: 1724–1730. F r catal gi g purp es only

156. Audronis, M., Leyland, A., Kelly, P. J., and Matthews,A. 2008. Composition and structure-property relationshipsof chromium-diboride/molybdenum-disulphide PVDnanocomposite hard coatings deposited by pulsed magnetronsputtering. Appl. Phys. A 91: 77–86.

157. Zabinski, J. S., Donley, M. S., and McDevitt, N. T.1993. Mechanistic study of the synergism between Sb 2 0 3and MoS 2 lubricant systems using Raman spectroscopy. Wear165: 103–108.

158. Scharf, T. W., Kotula, P. G., and Prasad, S. V. 2010.Friction and wear mechanisms in MoS 2 /Sb 2 O 3 /Au

nanocomposite coatings. Acta Mater. 58: 4100–4109.

159. Zabinski, J. S., Florkey, J. E., Walck, S. D.,Bultman, J. E., and McDevitt, N. T. 1995. Frictionproperties of WS 2 /graphite �uoride thin �lms grown bypulsed laser deposition. Surf. Coat. Technol. 76–77:400–406.

160. Stupp, B. C. 1981. Synergistic effects of metalsco-sputtered with MoS 2 . Thin Solid Films 84: 257–266.

161. Mulligan, C. P. and Gall, D. 2005. CrN-Agself-lubricating hard coatings. Surf. Coat. Technol. 200:1495–1500.

162. Mulligan, C. P., Blanchet, T. A., and Gall, D. 2010.CrN-Ag nanocomposite coatings: High-temperaturetribological response. Wear 269: 125–131.

163. Aouadi, S. M., Bohnhoff, A., Sodergen, M. et al. 2006.Tribological investigation of zirconium nitride/ silvernanocomposite structures. Surf. Coat. Technol. 201:418–422.

164. Köstenbauer, H., Fontalvo, G. A., Mitterer, C., andKeckes, J. 2008. Tribological properties of TiN/Agnanocomposite coatings. Tribol. Lett. 30: 53–60.

165. Ding, X.-Z., Zeng, X. T., and Goto, T. 2005.Unbalanced magnetron sputtered Ti-Si-N: MoS x compositecoatings for improvement of tribological properties. Surf.Coat. Technol. 198: 432–436.

166. Guleryuz, C. G., Krzanowski, J. E. Veldhuis, S. C.,and Fox-Rabinovich, G. S. 2009. Machining performance ofTiN coatings incorporating indium as a solid lubricant.Surf. Coat. Technol. 203: 3370–3376.

167. Krzanowski, J. E., Endrino, J. L., Nainaparambil, J.J., and Zabinski, J. S. 2003. Composite coatingsincorporating solid lubricant phases. Proc. 22nd HeatTreating Society Conference and the 2nd InternationalSurface Engineering Congress, Indiana, USA.

168. Basnyat, P., Luster, B., Kertzman, Z. et al. 2007.Mechanical and tribological properties of CrAlN-Agself-lubricating �lms. Surf. Coat. Technol. 202: 1011–1016.

169. Endrino, J. L., Nainaparampil, J. J., and Krzanowski,J. E. 2002. Magnetron sputter deposition of WC-Ag and

TiC-Ag coatings and their frictional properties in vacuumenvironments. Scr. Mater. 47: 613–618.

170. Suszko, T., Gulbiński, W., and Jagielski, J. 2006. Mo2 N/Cu thin �lms—The structure, mechanical andtribological properties. Surf. Coat. Technol. 200:6288–6292.

171. Mulligan, C. P., Blanchet, T. A., and Gall, D. 2010.Control of lubricant transport by a CrN diffusion barrierlayer during high-temperature sliding of a CrN-Ag compositecoating. Surf. Coat. Technol. 205: 1350–1355.

172. Srivastav, A., Kapoor, A., and Pathak, J. P. 1992. Therole of MoS 2 in hard overlay coatings of A1 2 O 3 indry sliding. Wear 155: 229–236.

173. Rahman, M., Haider, J., Dowling, D. P., Duggan, P.,and Hashmi, M. S. J. 2005. Investigation of mechanicalproperties of TiN + MoS x coating on plasma-nitridedsubstrate. Surf. Coat. Technol. 200: 1451–1457.

174. Guruvenket, S., Li, D., Klemberg-Sapieha, J. E.,Martinu, L., and Szpunar, J. 2009. Mechanical andtribological properties of duplex treated TiN, nc-TiN/a-SiNx and nc-TiCN/a-SiCN coatings deposited on 410 low alloystainless steel. Surf. Coat. Technol. 203: 2905–2911.

175. Barshilia, H. C., Deepthi, B., Srinivas, G., andRajam, K. S. 2012. Sputter deposited low-friction andtough Cr-Si 3 N 4 nanocomposite coatings on plasmanitrided M2 steel. Vacuum 86: 1118–1125.

176. Ochoa, E. A. and Figueroa, C. A. 2008. In�uence of themicrostructure on steel hardening in pulsed plasmanitriding. J. Vac. Sci. Technol. A 26: 328–332.

177. Wang, J., Xiong, J., Peng, Q. et al. 2009. Effects ofDC plasma nitriding parameters on microstructure andproperties of 304 L stainless steel. Mater. Character. 60:197–203.

178. Carrera, S., Salas, O., Moore, J. J., Woolverton, A.,and Sutter, E. 2003. Performance of CrN/MoS 2 (Ti)coatings for high wear low friction applications. Surf.Coat. Technol. 167: 25–32.

179. Dimigen, H. and Klages, C.-P. 1991. Microstructure andwear behavior of metal-containing diamond-like coatings.Surf. Coat. Technol. 49: 543–547.

180. Musil, J., Louda, M., Soukup, Z., and Kubásek, M.2008. Relationship between mechanical properties andcoef�cient of friction of sputtered a-C/Cu composite thin�lms. Diamond Relat. Mater. 17: 1905–1911.

181. Stüber, M., Leiste, H., Ulrich, S., Holleck, H., andSchild, D. 2002. Microstructure and properties of lowfriction TiC-C nanocomposite coatings deposited bymagnetron sputtering. Surf. Coat. Technol. 150: 218–226.

182. Gulbiński, W., Mathur, S., Shen, H., Suszko, T.,Gilewicz, A., and Warcholiński, B. 2005. Evaluation ofphase, composition, microstructure and properties inTiC/a-C:H thin �lms deposited by magnetron sputtering.Appl. Surf. Sci. 239: 302–310.

183. Lindquist, M., Wilhelmsson, O., Jansson, U., andWiklund, U. 2009. Tribo�lm formation and tribologicalproperties of TiC and nanocomposite TiAlC coatings. Wear266: 379–387.

184. Zheng, X. H., Tu, J. P., Gu, B., and Hu, S. B. 2008.Preparation and tribological behavior of TiN/a-C composite�lms deposited by DC magnetron sputtering. Wear 265:261–265.

185. Donnet, C. 1998. Recent progress on the tribology ofdoped diamond-like and carbon alloy coatings: A review.Surf. Coat. Technol. 100–101: 180–186.

186. Pauleau, Y. and Thièry, F. 2004. Deposition andcharacterization of nanostructured metal/carbon composite�lms. Surf. Coat. Technol. 180–181: 313–322.

187. Choi, J.-H., Ahn, H.-S., Lee, S.-C., and Lee, K.-R.2006. Stress reduction behavior in metal- incorporatedamorphous carbon �lms: First-principle approach. J. Phys.Conf. Series 29: 155–158.

188. Corbella. C., Bertran, E., Polo, M. C., Pascual, E.,and Andújar, J. L. 2007. Structural effects ofnanocomposite �lms of amorphous carbon and metal depositedby pulsed-DC reactive magnetron sputtering. Diamond Relat.Mater. 16: 1828–1834.

189. Matenoglou, G., Evangelakis, G. A., Kosmidis, C.,Foulias, S., Papadimitriou, D., and Patsalas, P. 2007.Pulsed laser deposition of amorphous carbon/silvernanocomposites. Appl. Surf. Sci. 253: 8155–8159.

190. Grill, A. 1999. Diamond-like carbon: State of the art.Diamond Relat. Mater. 8: 428–434.

191. Doll, G. L. 1999. Deposition, characterization, andapplications of metal-doped diamond-like carbon �lms. InSurface Engineering: Science and Technology I, ed. A.Kumar, Y.-W. Chung, J. J. Moore and J. E. Smugeresky, pp.295–305. Pennsylvania: The Minerals, Metals & MaterialsSociety.

192. Wu, J.-H., Sanghavi, M., Sanders, J. H., Voevodin, A.A., Zabinski, J. S., and Rigney, D. A. 2003. Slidingbehavior of multifunctional composite coatings based ondiamond-like carbon. Wear 255: 859–868.

193. Kao, W. H., Su, Y. L., and Yao, S. H. 2006.Tribological property and drilling application of Ti-C:Hand Cr-C:H coatings on high-speed steel substrates. Vacuum80: 604–614.

194. Tan, M., Zhu, J., Han, J., Han, X., Niu, L., and Chen,W. 2007. Stress evolution of tetrahedral amorphous carbonupon boron incorporation. Scripta Mater. 57: 141–144.

195. Sánchez-López, J. C., and Fernández, A. 2008. Dopingand alloying effects on DLC coatings. In Tribology ofDiamond-Like Carbon Coatings: Fundamentals andApplications, ed. C. Donnet and A. Erdemir, pp. 311–338.New York: Springer.

196. Harris, S. J., Weiner, A. M., and Meng, W.-J. 1997.Tribology of metal-containing diamond-like carboncoatings. Wear 211: 208–217.

197. Wang, D.-Y., Weng, K.-W., Chang, C.-L., and Guo, X.-J.2000. Tribological performance of metal doped diamond-likecarbon �lms deposited by cathodic arc evaporation. DiamondRelat. Mater. 9: 831–837.

198. Gilmore, R. and Hauert, R. 2000. Comparative study ofthe tribological moisture sensitivity of Si-free andSi-containing diamond-like carbon �lms. Surf. Coat.Technol. 133–134: 437–442.

199. Miyake, S. and Kaneko, R. 1992. Microtribologicalproperties and potential applications of hard, lubricatingcoatings. Thin Solid Films 212: 256–261.

200. Gilmore, R. and Hauert, R. 2001. Control of the

tribological moisture sensitivity of diamond-like carbon�lms by alloying with F, Ti or Si. Thin Solid Films398–399: 199–204.

201. Oguri, K. and Arai, T. 1990. Low friction coatings ofdiamond-like carbon with silicon prepared byplasma-assisted chemical vapor deposition. J. Mater. Res.5: 2567–2571.

202. Pauleau, Y., Thièry, F., Uglov, V. V., Kuleshov, A.K., Dub, S. N., and Samtsov, M. P. 2003. Preparation,structure and mechanical properties of nanostructuredcopper-carbon �lms. Rev. Adv. Mater. Sci. 4: 139–146.

203. Martínez-Martínez, D., López-Cartes, C., Justo, A.et al. 2005. Tailored synthesis of TiC/a-C nanocompositetribological coatings. J. Vac. Sci. Technol. A 23:1732–1736.

204. Voevodin, A. A., Prasad, S. V., and Zabinski, J. S.1997. Nanocrystalline carbide/amorphous carbon composites.J. Appl. Phys. 82: 855–858.

205. Voevodin, A. A., O’Neill, J. P., and Zabinski, J. S.1999. Tribological performance and tribochemistry ofnanocrystalline WC/amorphous diamond-like carboncomposites. Thin Solid Films 342: 194–200.

206. Gassner, G., Patscheider, J., Mayrhofer, P. H., Šturm,S., Scheu, C., and Mitterer, C. 2007. Tribologicalproperties of nanocomposite CrC x /a-C:H thin �lms. Tribol.Lett. 27: 97–104. For cataloging p rposes nly

207. Noshiro, J., Watanabe, S., Sakurai, T., and Miyake, S.2006. Friction properties of co-sputtered sul�de/ DLC solidlubricating �lms. Surf. Coat. Technol. 200: 5849–5854.

208. Dorfman, V. F. 1992. Diamond-like nanocomposites(DLN). Thin Solid Films 212: 267–273.

209. Outten, C. A., Kester, D., Venkatraman, C., Bray, D.,and Halter, C. Diamond-like nanocomposite coatings forspace applications. In Surface Engineering: Science andTechnology I, ed. A. Kumar, Y.-W. Chung, J. J. Moore andJ. E. Smugeresky, pp. 271–282. Pennsylvania: The Minerals,Metals & Materials Society.

210. Neerinck, D. and Goel, A. 2001. Diamond-LikeNanocomposite Compositions. United States Patent No.6,200,675 B1.

211. Neerinck, D., Persoone, P., Sercu, M., Goel, A.,Kester, D., and Bray, D. 1998. Diamond-like nanocompositecoatings (a-C:H/a-Si:O) for tribological applications.Diamond Relat. Mater. 7: 468–471.

212. Neerinck, D., Persoone, P., Sercu, M. et al. 1998.Diamond-like nanocomposite coatings for low-wear andlow-friction applications in humid environments. Thin SolidFilms 317: 402–404.

213. Kester, D. J., Brodbeck, C. L., Singer, I. L., andKyriakopoulos, A. 1999. Sliding wear behavior ofdiamond-like nanocomposite coatings. Surf. Coat. Technol.113: 268–273.

214. Scharf, T. W., Ohlhausen, J. A., Tallant, D. R., andPrasad, S. V. 2007. Mechanisms of friction in diamond likenanocomposite coatings. J. Appl. Phys. 101: 063521 1–11.

215. Adachi, K. and Kato, K. 2008. Tribology of carbonnitride coatings. In Tribology of Diamond-Like CarbonFilms: Fundamentals and Applications, ed. C. Donnet and A.Erdemir, pp. 339–361. New York: Springer.

216. Sánchez-López, J. C., Belin, M., Donnet, C., Quirós,C., and Elizalde, E. 2002. Friction mechanisms ofamorphous carbon nitride �lms under variable environments:A triboscopic study. Surf. Coat. Technol. 160: 138–144.

217. Broitman, E., Neidhardt, J., and Hultman, L. 2008.Fullerene-like carbon nitride: A new carbon-basedtribological coating. In Tribology of Diamond-Like CarbonFilms: Fundamentals and Applications, ed. C. Donnet and A.Erdemir, pp. 620–653. New York: Springer.

218. Grimanelis, D., Yang, S., Böhme, O. et al. 2002.Carbon based coatings for high temperature cutting toolapplications. Diamond Relat. Mater. 11: 176–184.

219. Park, I.-W., Mishra, B., Kim, K. H., and Moore, J. J.2007. Multifunctional Ti-Si-B-C-N tribologicalnanocomposite coatings for aerospace applications. Mater.Sci. Forum 539–543: 173–180.

220. Gupta, S., Filimonov, D., Palanisamy, T., El-Raghy,T., and Barsoum, M. W. 2007Ta 2 AlC and Cr 2 AlC Ag-basedcomposites-New solid lubricant materials for use over awide temperature range against Ni-based superalloys andalumina. Wear 262: 1479–1489.

221. Zhang, S., Fu, Y., Du, H., Zeng, X. T., and Liu, Y. C.2002. Magnetron sputtering of nanocomposite (Ti,Cr)CN/DLCcoatings. Surf. Coat. Technol. 162: 42–48.

222. Sliney, H. E. 1982. Solid lubricant materials for hightemperatures—A review. Tribol. Int. 15: 303–315.

223. Gulbiński, W. and Suszko, T. 2006. Thin �lms of MoO 3-Ag 2 O binary oxides—The high temperature lubricants.Wear 261: 867–873.

224. Ouyang, J. H., Sasaki, S., and Umeda, K. 2001.Low-pressure plasma-sprayed ZrO 2 -CaF 2 compositecoating for high temperature tribological applications.Surf. Coat. Technol. 137: 21–30.

225. Muratore, C. and Voevodin, A. A. 2009. Chameleoncoatings: Adaptive surfaces to reduce friction and wear inextreme environments. Annu. Rev. Mater. Res. 39: 297–324.

226. Aouadi, S. M., Luster, B., Kohli, P., Muratore, C.,and Voevodin, A. A. 2009. Progress in the development ofadaptive nitride-based coatings for high temperaturetribological applications. Surf. Coat. Technol. 204:962–968.

227. Voevodin, A. A., O’Neill, J. P., and Zabinski, J. S.1999. Nanocomposite tribological coatings for aerospaceapplications. Surf. Coat. Technol. 116–119: 36–45.

228. Voevodin, A. A., and Zabinski, J. S. 2005.Nanocomposite and nanostructured tribological materials forspace applications. Compos. Sci. Technol. 65: 741–748.

229. Voevodin, A. A. and Zabinski, J. S. 2000. Supertoughwear-resistant coatings with ‘chameleon’ surfaceadaptation. Thin Solid Films 370: 223–231.

230. Baker, C. C., Chromik, R. R., Wahl, K. J., Hu, J. J.,and Voevodin, A. A. 2007. Preparation of chameleoncoatings for space and ambient environments. Thin SolidFilms 515: 6737–3743.

231. Baker, C. C., Hu, J. J., Voevodin, A. A. et al. 2006.Preparation of Al 2 O 3 /DLC/Au/MoS 2 chameleon coatingsfor space and ambient environments. Surf. Coat. Technol.201: 4224–4229.

232. Muratore, C., Voevodin, A. A., Hu, J. J., and

Zabinski, J. S. 2006. Multilayered YSZ-Ag-Mo/TiN adaptivetribological nanocomposite coatings. Tribol. Lett. 24:201–206.

233. Muratore, C., Voevodin, A. A., Hu, J. J., andZabinski, J. S. 2006. Tribology of adaptive nanocompositeyttria-stabilized zirconia coatings containing silver andmolybdenum from 25 to 700°C. Wear 261: 797–805.

234. Hu, J. J., Muratore, C., and Voevodin, A. A. 2007.Silver diffusion and high-temperature lubricationmechanisms of YSZ-Ag-Mo based nanocomposite coatings.Compos. Sci. Technol. 67: 336–347.

235. Zabinski, J. S., Bultman, J. E., Sanders, J. H., andHu, J. J. 2006. Multi-environmental lubrication performanceand lubrication mechanism of MoS 2 /Sb 2 O 3 /C composite�lms. Tribol. Lett. 23: 155–163.

236. Aouadi, S. M., Singh, D. P., Stone, D. S. et al. 2010.Adaptive VN/Ag nanocomposite coatings with lubriciousbehavior from 25 to 1000°C. Acta Mater. 58: 5326–5331.

237. Pearson, J. D., Zikry, M. A., and Wahl, K. J. 2009.Microstructural modeling of adaptive nanocompositecoatings for durability and wear. Wear 266: 1003–1012.

238. Aouadi, S. M., Paudel, Y., Luster, B. et al. 2008.Adaptive Mo 2 N/MoS 2 /Ag tribological nanocompositecoatings for aerospace applications. Tribol. Lett. 29:95–103.

239. Prasad, S. V., McDevitt, N. T., and Zabinski, J. S.2000. Tribology of tungsten disul�de-nanocrystalline zincoxide adaptive lubricant �lms from ambient to 500°C. Wear237: 186–196.

240. Walck, S. D., Zabinski, J. S., McDevitt, N. T., andBultman, J. E. 1997. Characterization of air-annealed,pulsed laser deposited ZnO-WS 2 solid �lm lubricants bytransmission electron microscopy. Thin Solid Films 305:130–143.

241. Zabinski, J. S., Donley, M. S., Dyhouse, V. J., andMcDevitt, N. T. 1992. Chemical and tribologicalcharacterization of PbO-MoS 2 �lms grown by pulsed laserdeposition. Thin Solid Films 214: 156–163.

242. Gassner, G., Mayrhofer, P. H., Kutschej, K., Mitterer,C., and Kathrein, M. 2004. A new low friction concept for

high temperatures: Lubricious oxide formation on sputteredVN coatings. Tribol. Lett. 17: 751–756.

243. Fateh, N., Fontalvo, G. A., Gassner, G., and Mitterer,C. 2007. The bene�cial effect of high-temperatureoxidation on the tribological behaviour of V and VNCoatings. Tribol. Lett. 28: 1–7.

244. Fateh, N., Fontalvo, G. A., and Mitterer, C. 2008.Tribological properties of reactive magnetron sputtered V2 O 5 and VN-V 2 O 5 coatings. Tribol. Lett. 30: 21–26.

245. John, P. J., Prasad, S. V., Voevodin, A. A., andZabinski, J. S. 1998. Calcium sulphate as a hightemperature solid lubricant. Wear 219: 155–161.

246. Pauleau, Y., Juliet, P., and Gras. R. 1998.Tribological properties of calcium �uoride-based solidlubricant coatings at high temperatures. Thin Solid Films317: 481–485.

247. Voevodin, A. A., and Zabinski, J. S. 1998.Load-adaptive crystalline-amorphous nanocomposites.J. Mater. Sci. 33: 319–327.

248. Voevodin, A. A., Hu, J. J., Fitz, T. A., and Zabinski,J. S. 2001. Tribological properties of adaptivenanocomposite coatings made of yttria stabilized zirconiaand gold. Surf. Coat. Technol. 146–147: 351–356.

249. Voevodin, A. A., Fitz, T. A., Hu, J. J., and Zabinski,J. S. 2002. Nanocomposite tribological properties coatingswith “chameleon” surface adaptation. J. Vac. Sci. Technol.A 20:1434–1444.

250. Muratore, C., Hu, J. J., and Voevodin, A. A. 2009.Tribological coatings for lubrication over multiplethermal cycles. Surf. Coat. Technol. 203: 957–962.

251. Muratore, C., Clarke, D. R., Jones. J. G., andVoevodin, A. A. 2008. Smart tribological coatings with wearsensing capability. Wear 265: 913–920.

252. Stone, D., Liu, J., Singh, D. P. et al. 2010. Layeredatomic structures of double oxides for low shear strengthat high temperatures. Scripta Mater. 62: 735–738.

253. Muratore, C. and Voevodin, A. A. 2006. Molybdenumdisulphide as a lubricant and catalyst in adaptivenanocomposite coatings. Surf. Coat. Technol. 201:

4125–4130.

254. Tenne, R., Homyonfer, M., and Feldman, Y. 1998.Nanoparticles of layered compounds with hollow cagestructures (inorganic fullerene-like structures). Chem.Mater. 10: 3225–3238.

255. Rapoport, L., Bilik, Y., Feldman, Y., Homyonfer, M.,Cohen, S. R., and Tenne, R. 1997. Hollow nanoparticles ofWS 2 as potential solid-state lubricants. Nature 387:791–793.

256. Joly-Pottuz, L. and Dassenoy, F. 2008. Nanoparticlesmade of metal dichalcogenides. In Nanolubricants, ed.J.-M. Martin and N. Ohmae, pp. 15–92. England: John Wiley &Sons Ltd.

257. Zhu, Y. Q., Sekine, T., Li, Y. H. et al. 2005. WS 2and MoS 2 inorganic fullerenes—Super shock absorbers atvery high pressures. Adv. Mater. 17: 1500–1503.

258. Rapoport, L., Fleischer, N., and Tenne, R. 2005.Applications of WS 2 (MoS 2 ) inorganic nanotubes andfullerene-like nanoparticles for solid lubrication and forstructural nanocomposites. J. Mater. Chem. 15: 1782–1788.

259. Tenne, R. 2006. Inorganic nanotubes and fullerene-likenanoparticles. Nat. Nanotechnol. 1: 103–111. For cat loginpur os s nly

260. Yang, H., Liu, S., Li, J., Li, M., Peng, G., and Zou,G. 2006. Synthesis of inorganic fullerene-like WS 2nanoparticles and their lubricating performance.Nanotechnology 17: 1512–1519.

261. Tenne, R. and Redlich, M. 2010. Recent progress in theresearch of inorganic fullerene-like nanoparticles andinorganic nanotubes. Chem. Soc. Rev. 39: 1423–1434.

262. Chhowalla, M. and Amaratunga, G. A. J. 2000. Thin �lmsof fullerene-like MoS 2 nanoparticles with ultralowfriction and wear. Nature 407: 164–167.

263. Tenne, R. and Seifert, G. 2009. Recent progress in thestudy of inorganic nanotubes and fullerene-likestructures. Annu. Rev. Mater. Res. 39: 387–413.

264. Tenne, R., Rapoport, L., Lvovsky, M., Feldman, Y., andLeshchinsky, V. 2003. Hollow fullerene-like nanoparticlesas solid lubricants in composite metal matrices. United

States Patent No. 2003/ 0144155 A1.

265. Tenne, R., Feldman, Y., Alla, Z., and Rita, R. 2001.Inorganic fullerene-like tungsten disulphide hollownanoparticles and nanotubes. European Patent No. EP 1 840088 A1.

266. Zou, T., Tu, J., Huang, H., Lai, D., Zhang, L., andHe, D. 2006. Preparation and tribological properties ofinorganic fullerene-like MoS 2 . Adv. Eng. Mater. 8:289–293.

267. Hu, J. J. and Zabinski, J. S. 2005. Nanotribology andlubrication mechanisms of inorganic fullerene-like MoS 2nanoparticles investigated using lateral force microscopy(LFM). Tribol. Lett. 18: 173–180.

268. Hu, J. J., Bultman, J. E., and Zabinski, J. S. 2004.Inorganic fullerene-like nanoparticles produced by arcdischarge in water with potential lubricating ability.Tribol. Lett. 17: 543–546.

269. José-Yacamán, M., López, H., Santiago, P., Galván, D.H., Garzón, I. L., and Reyes, A. 1996. Studies of MoS 2structures produced by electron irradiation. Appl. Phys.Lett. 69: 1065–1067.

270. Sen, R., Govindaraj, A., Suenaga, K. et al. 2001.Encapsulated and hollow closed-cage structures of WS 2 andMoS 2 prepared by laser ablation at 450–1050°C. Chem.Phys. Lett. 340: 242–248.

271. Rapoport, L., Fleischer, N., and Tenne, R. 2003.Fullerene-like WS 2 nanoparticles: Superior lubricantsfor harsh conditions. Adv. Mater. 15: 651–655.

272. Drummond, C., Alcantar, N., Israelachvili, J., Tenne,R., and Golan, Y. 2001. Microtribology and friction-inducedmaterial transfer in WS 2 nanoparticle additives. Adv.Funct. Mater. 11: 348–354.

273. Rapoport, L., Feldman, Y., Homyonfer, M. et al. 1999.Inorganic fullerene-like material as additives tolubricants: Structure-function relationship. Wear 225–229:975–982.

274. Hou, X., Shan, C. X., and Choy, K. -L. 2008.Microstructures and tribological properties of PEEK-basednanocomposite coatings incorporating inorganicfullerene-like nanoparticles. Surf. Coat. Technol. 202:

2287–2291.

275. Greenberg, R., Halperin, G., Etsion, I., and Tenne, R.The effect of WS 2 nanoparticles on friction reduction invarious lubrication regimes. Tribol. Lett. 17: 179–186.

276. Rapoport, L., Lvovsky, M., Lapsker, I. et al. 2001.Slow release of fullerene-like WS 2 nanoparticles fromFe-Ni graphite matrix: A self-lubricating nanocomposite.Nano Lett. 1: 137–140.

277. Hübert, T., Hattermann, H., and Griepentrog, M. 2009.Sol-gel-derived nanocomposite coatings �lled withinorganic fullerene-like WS 2 . J. Sol-Gel Sci. Technol.51: 295–300.

278. Cizaire, L., Vacher, B., Mogne, T. L. et al. 2002.Mechanisms of ultra-low friction by hollow inorganicfullerene-like MoS 2 nanoparticles. Surf. Coat. Technol.160: 282–287.

279. Rapoport, L., Nepomnyashchy, O., Verdyan, A. et al.2004. Polymer nanocomposites with fullerene-like solidlubricant. Adv. Eng. Mater. 6: 44–48.

280. Bahadur, S. and Gong, D. 1992. The action of �llers inthe modi�cation of the tribological behavior of polymers.Wear 158: 41–59.

281. Fusaro, R. L. 1990. Self-lubricating polymercomposites and polymer transfer �lm lubrication for spaceapplications. Tribol. Int. 23: 105–122.

282. Sawyer, W. G., Freudenberg, K. D., Bhimaraj, P., andSchadler, L. S. 2003. A study on the friction and wearbehavior of PTFE �lled with alumina nanoparticles. Wear254: 573–580.

283. Khedkar, J., Negulescu, I., and Meletis, E. I. 2002.Sliding wear behavior of PTFE composites. Wear 252:361–369.

284. McCook, N. L., Boesl, B., Burris, D. L., and Sawyer,W. G. 2006. Epoxy, ZnO, and PTFE nanocomposite: Frictionand wear optimization. Tribol. Lett. 22: 253–257.

285. Blanchet, T. A. and Kennedy, F. E. 1992. Sliding wearmechanism of polytetra�uoroethylene (PTFE) and PTFEcomposites. Wear 153: 229–243.

286. Chen, W. X., Li, F., Han, G. et al. 2003. Tribologicalbehavior of carbon-nanotube-�lled PTFE composites. Tribol.Lett. 15: 275–278.

287. McCook, N. L., Burris, D. L., Bourne, G. R., Steffens,J., Hanrahan, J. R., and Sawyer, W. G. 2005. Wearresistant solid lubricant coating made from PTFE and epoxy.Tribol. Lett. 18: 119–124.

288. McCook, N. L., Hamilton, M. A., Burris, D. L., andSawyer, W. G. 2007. Tribological results of PEEKnanocomposites in dry sliding against 440°C in various gasenvironments. Wear 262: 1511–1515.

289. Shi, Y., Mu, L., Feng, X., and Lu, X. 2011. Thetribological behavior of nanometer and micrometer TiO 2particle-�lled polytetra�uoroethylene/polyimide. Mater.Des. 32: 964–970.

290. Xiang, D. and Shan, K. 2006. Friction and wearbehavior of self-lubricating and heavily loaded metalPTFEcomposites. Wear 260: 1112–1118.

291. Zhang, X.-R., Pei, X.-Q., and Wang, Q.-H. 2009.Friction and wear studies of polyimide composites �lledwith short carbon �bers and graphite and micro SiO 2 .Mater. Des. 30: 4414–4420.

292. Bahadur, S. and Gong, D. 1992. The role of coppercompounds as �llers in the transfer and wear behavior ofpolyetheretherketone. Wear 154: 151–165.

293. Li, F., Hu, K.-A., Li, J.-L., and Zhao, B.-Y. 2002.The friction and wear characteristics of nanometer ZnO�lled polytetra�uoroethylene. Wear 249: 877–882.

294. Friedrich, K., Zhang, Z., and Schlarb, A. K. 2005.Effects of various �llers on the sliding wear of polymercomposites. Compos. Sci. Technol. 65: 2329–2343.

295. Gong, D., Zhang, B., Xue, Q.-J., and Wang, H.-L. 1990.Investigation of adhesion wear of �lledpolytetra�uoroethylene by ESCA, AES and XRD. Wear 57:25–39.

296. Bahadur, S. and Tabor, D. 1984. The wear of �lledpolytetra�uoroethylene. Wear 98: 1–13.

297. Briscoe, B. J. and Sinha, S. K. Tribologicalapplications of polymers and their composites: Past,

present and future prospects. 2008. In Tribology ofpolymeric Nanocomposites, ed. K. Friedrich andA. K. Schlarb, pp. 1–14. Oxford: Elsevier.

298. Yin, G. L., Huang, P. H., Yu, Z., He, D. N., and Tu,J. P. 2006. Microstructure, chemical and tribologicalinvestigations of Mo x W 1-x S y co-sputtered composite�lms. Tribol. Lett. 22: 37–43.

299. Uglova, V. V., Kuleshov, A. K., Astashynskaya, M. V.et al. 2005. Mechanical properties of copper/carbonnanocomposite �lms formed by microwave plasma assisteddeposition techniques from argon–methane andargon–acetylene gas mixtures. Compos. Sci. Technol. 65:785–791.

300. Wilhelmsson, O., Råsander, M., Carlsson, M. et al.2007. Design of nanocomposite low-friction coatings. Adv.Funct. Mater. 17: 1611–1616.

301. Shaha, K. P., Pei, Y. T., Chen, C. Q., and De Hosson,J. T. M. 2010. Synthesis of ultra-smooth and ultralowfriction DLC based nanocomposite �lms on rough substrates.Thin Solid Films 519: 1618–1622.

302. Arai, S., Fujimori, A., Murai, M., and Endo, M. 2008.Excellent solid lubrication of electrodepositednickel-multiwalled carbon nanotube composite �lms. Mater.Lett. 62: 3545–3548.

303. Cardinal, M. F., Castro, P. A., Baxi, J., Liang, H.,and Williams, F. J. 2009. Characterization and frictionalbehavior of nanostructured Ni-W-MoS 2 composite coatings.Surf. Coat. Technol. 204: 85–90.

304. Nickchi, T., Ghorbani, M., Alfantazi, A., and Farhat,Z. 2011. Fabrication of low friction bronze-graphitenano-composite coatings. Mater. Des. 32: 3548–3553.

305. Liang, X.-S., Ouyang, J.-H., Li, Y.-F., and Wang,Y.-M. 2009. Electrodeposition and tribological propertiesof Ni-SrSO 4 composite coatings. Appl. Surf. Sci. 255:4316–4321.

306. Sun, X. J. and Li, J. G. 2007. Friction and wearproperties of electrodeposited nickel-titania nanocompositecoatings. Tribol. Lett. 28: 223–228.

307. Ploof, L. 2008. Electroless nickel composite coatings.Adv. Mater. Process. 166: 36–38.

308. Zhang, Y. Z., Wu, Y. Y., Sun, K. N., and Yao, M. 1998.Characterization of electroless Ni-P-PTFE compositedeposits. J. Mater. Sci. Lett. 17: 119–122.

309. Ramalho, A. and Miranda, J. C. 2005. Friction and wearof electroless NiP and NiP + PTFE coatings. Wear 259:828–834.

310. Sahoo, P. and Das, S. K. 2011. Tribology ofelectroless nickel coatings—A review Mater. Des. 32:1760–1775.

311. Wu, Y., Liu, H., Shen, B., Liu, L., and Hu, W. 2006.The friction and wear of electroless Ni-P matrix with PTFEand/or SiC particles composite. Tribol. Int. 39: 553–559.

312. Sarret, M., Müller, C., and Amell, A. 2006.Electroless NiP micro- and nano-composite coatings. Surf.Coat. Technol. 201: 389–395.

313. Feldstein, M. D. 1998. Composite electroless nickelcoatings for the aerospace and airline industries. Plat.Surf. Finish. 85: 248–252.

314. Grosjean, A., Rezrazi, M., Takadoum, J., and Berçot,P. 2001. Hardness, friction and wear characteristics ofnickel-SiC electroless composite deposits. Surf. Coat.Technol. 137: 92–96.

315. Chen, X. H., Chen, C. S., Xiao, H. N. et al. 2006. Dryfriction and wear characteristics of nickel/carbonnanotube electroless composite deposits. Tribol. Int. 39:22–28.

316. Taylor, D. J., Fleig, P. F., Schwab, S. T., and Page,R. A. 1999. Sol-gel derived, nanostructured oxidelubricant coatings. Surf. Coat. Technol. 120–121: 465–469.For cat loging pu pos s only

317. Hai-Dou, W., Da-Ming, Z., Kun-Lin, W., Jia-Jun, L.,Yue, C., and Zhou-Ping, C. 2003. Study on tribologicalproperties of iron sul�de coatings prepared by a sol-gelmethod. J. Mater. Sci. Lett. 22: 1043–1046.

318. Taylor, D. J., Fleig, P. F., and Page, R. A. 2002.Characterization of nickel titanate synthesized by sol-gelprocessing. Thin Solid Films 408: 104–110.

319. Skeldon, P., Wang, H. W., and Thompson, G. E. 1997.

Formation and characterization of self-lubricating MoS 2precursor �lms on anodized aluminium. Wear 206: 187–196.

320. Wang, H. and Wang, H. 2007. Fabrication ofself-lubricating coating on aluminum and its frictionalbehavior. Appl. Surf. Sci. 253: 4386–4389.

321. Spalvins, T. 1978. Coatings for wear and lubrication.Thin Solid Films 53: 285–300.

322. Zabinski, J. S., Voevodin, A. A., Nainaparambil, J.J., Prasad, S. V., and Pierce, N. A. 1999. Novel thin �lmsolid lubricants and hard coatings by pulsed laser basedtechniques. In Surface Engineering: Science and TechnologyI, ed. A. Kumar, Y. -W. Chung, J. J. Moore and J. E.Smugeresky, pp. 127–142. Pennsylvania: The Minerals,Metals & Materials Society.

323. Wang, H.-D., Xu, B.-S., Liu, J.-J., and Zhuang, D.-M.2005. Characterization and tribological properties ofplasma sprayed FeS solid lubrication coatings. Mater.Charact. 55: 43–49.

324. Sliney, H. E. 1979. Wide temperature spectrumself-lubricating coatings prepared by plasma spraying.Thin Solid Films 64: 211–217.

325. Chrisey, D. B. and Hubler, G. K. 1994. Pulsed LaserDeposition of Thin Films. New York: John Wiley.

326. Rigato, V., Maggioni, G., Patelli, A., Boscarino, D.,Renevier, N. M., and Teer, D. G. 2000. Properties ofsputter-deposited MoS 2 -metal composite coatings depositedby closed �eld unbalanced magnetron sputter ion plating.Surf. Coat. Technol. 131: 206–210.

327. Voevodin, A. A., Capano, M. A., Safriet, A. J.,Donley, M. S., and Zabinski, J. S. 1996. Combined magnetronsputtering and pulsed laser deposition of carbides anddiamond-like carbon �lms. Appl. Phys. Lett. 69: 188–190.

328. Buck, V. 1987. Preparation and properties of differenttypes of sputtered MoS 2 �lms. Wear 114: 263–274.

329. Weiß, V., Seeger, S., Ellmer, K., and Mientus, R.2007. Reactive magnetron sputtering of tungsten disul�de(WS 2-x ) �lms: In�uence of deposition parameters ontexture, microstructure, and stoichiometry. J. Appl. Phys.101: 103502 1–9.

330. Dimigen, H., Hübsch, H., Willich, P., and Reichelt, K.1985. Stoichiometry and friction properties of sputteredMoS x layers. Thin Solid Films 129: 79–91.

331. Bichsel, R., Buffat, P., and Levy, F. 1986.Correlation between process conditions, chemicalcomposition and morphology of MoS 2 �lms prepared by RFplanar magnetron sputtering. J. Phys. D.: Appl. Phys. 19:1575–1585.

332. Buck. V. 1986. A neglected parameter (watercontamination) in sputtering of MoS 2 �lms. Thin SolidFilms 139: 157–168.

333. Lince, J. R. and Fleischauer, P. D. 1987.Crystallinity of rf-sputtered MoS 2 �lms. J. Mater. Res.2: 827–838.

334. Spalvins, T. 1980. Tribological properties ofsputtered MoS 2 �lms in relation to �lm morphology. ThinSolid Films 73: 291–297.

335. Simmonds, M. C., Savan, A., P�üger, E., andSwygenhoven, H. V. 2001. Microstructure and tribologicalperformance of MoS x /Au co-sputtered composites. J. Vac.Sci. Technol. A 19: 609–613.

336. Fleischauer, P. D., Lince, J. R., Bertrand, P. A., andBauer, R. 1989. Electronic structure and lubricationproperties of MoS 2 : A qualitative molecular orbitalapproach. Langmuir 5: 1009–1015.

337. Charitidis, C. A. 2010. Nanomechanical andnanotribological properties of carbon-based thin �lms: Areview. Int. J. Refract. Met. Hard Mater. 28: 51–70.

338. Corbella, C., Pascual, E., Oncins, G., Canal, C.,Andújar, J. L., and Bertran, E. 2005. Composition andmorphology of metal-containing diamond-like carbon �lmsobtained by reactive magnetron sputtering. Thin SolidFilms 482: 293–298.

339. Corbella, C., Oncins, G., Gómez, M. A. et al. 2005.Structure of diamond-like carbon �lms containingtransition metals deposited by reactive magnetronsputtering. Diamond Relat. Mater. 14: 1103–1107.

340. Irmer, G. and Dorner-Reisel, A. 2005. Micro-Ramanstudies on DLC coatings. Adv. Eng. Mater. 7: 694–705.

341. Zhang, S., Jen, T. M., Zeng, X., Xie, H., and Hing, P.2000. Raman and PEELS studies of magnetron sputtered a-C.Int. J. Mod. Phys. B 14: 268–273.

342. Tay, B. K., Shi, X., Tan, H. S., Yang, H. S., and Sun,Z. 1998. Raman studies of tetrahedral amorphous carbon�lms deposited by �ltered cathodic vacuum arc. Surf. Coat.Technol. 105: 155–158.

343. Cui, W. G., Lai, Q. B., Zhang, L., and Wang, F. M.2010. Quantitative measurements of sp 3 content in DLC�lms with Raman spectroscopy. Surf. Coat. Technol. 205:1995–1999.

344. Ferrari, A. C. and Robertson, J. 2000. Interpretationof Raman spectra of disordered and amorphous carbon. Phys.Rev. B 61: 14095–14107.

345. Prawer, S., Nugent, K. W., Lifshitz, Y. et al. 1996.Systematic variation of the Raman spectra of DLC �lms as afunction of sp 2 :sp 3 composition. Diamond Relat. Mater.5: 433–438.

346. Ferrari, A. C. 2008. Non-destructive characterizationof carbon �lms. In Tribology of Diamond-Like Carbon Films:Fundamentals and Applications, ed. C. Donnet and A.Erdemir, pp. 25–82. New York: Springer.

347. LiBassi, A., Ferrari, A. C., Stolojan, V., Tanner, B.K., Robertson, J., and Brown, L. M. 2000. Density, sp 3content and internal layering of DLC �lms by X-rayre�ectivity and electron energy loss spectroscopy. DiamondRelat. Mater. 9: 771–776.

348. Silva, S. R. P. and Stolojan, V. 2005. Electron energyloss spectroscopy of carbonaceous materials. Thin SolidFilms 488: 283–290.

349. Maydell, E. A., Dunlop, E., Fabian, D. J., Haupt, J.,and Gissler, W. 1993. Electron energy loss study ofdiamond-like and amorphous carbon �lms. Diamond Relat.Mater. 2: 873–878.

350. Ferrari, A. C., Kleinsorge, B., Adamopoulos, G. et al.2000. Determination of bonding in amorphous carbons byelectron energy loss spectroscopy, Raman scattering andX-ray re�ectivity. J. Non-Cryst. Solids. 266–269: 765–768.

351. Kovarik, P., Bourdon, E. B. D., and Prince, R. H.1993. Electron-energy-loss characterization of

laserdeposited a-C, a-C:H, and diamond �lms. Phys. Rev. B48: 12123–12129.

352. Bhushan, B. 2008. Nanotribology of ultrathin and hardamorphous carbon �lms. In Tribology of Diamond-Like CarbonFilms: Fundamentals and Applications, ed. C. Donnet and A.Erdemir, pp. 510–570. New York: Springer.

353. Ni, W., Cheng, Y.-T., Lukitsch, M. J., Weiner, A. M.,Lev, L. C., and Grummon, D. S. 2004. Effects of the ratioof hardness to Young’s modulus on the friction and wearbehavior of bilayer coatings. Appl. Phys. Lett. 85:4028–4030.

354. Leyland, A. and Matthews. A. 2000. On the signi�canceof the H/E ratio in wear control: A nanocomposite coatingapproach to optimized tribological behavior. Wear 246:1–11.

355. Shtanskiı˘, D. V., Kulinich, S. A., Levashov, E. A.,and Moore, J. J. 2003. Structure and physical-mechanicalproperties of nanostructured thin �lms. Phys. Solid State45: 1177–1184.

356. Beake, B. D., Vishnyakov, V. M., Valizadeh, R., andColligon, J. S. 2006. In�uence of mechanical properties onthe nanoscratch behaviour of hard nanocomposite TiN/Si 3 N4 coatings on Si. J. Phys. D Appl. Phys. 39: 1392–1397.

357. Zhang, S., Sun, D., and Bui, X. L. 2007. Magnetronsputtered hard and yet tough nanocomposite coatings withcase studies: Nanocrystalline TiN embedded in amorphous SiNx . In Nanocomposite Thin Films and Coatings, ed. S. Zhangand N. Ali, pp. 1–108 London: Imperial College Press.

358. Brown, R. 1970. Thin �lm substrates. In Handbook ofThin Film Technology, ed. L. I. Maissel and R. Glang, pp.6 1–50. New York: McGraw-Hill, Inc.

359. Thornton, J. A. and Penfold, A. S. 1978. Cylindricalmagnetron sputtering. In Thin Film Processes, ed. J. L.Vossen and W. Kern, pp. 75–113. New York: Academic Press,Inc.

360. Budtz-Jørgensen, C. V., Kringhøj, P., and Bøttiger, J.1999. The critical role of hydrogen for physicalsputtering with Ar-H 2 glow discharges. Surf. Coat.Technol. 116–119: 938–943.

361. Ji, L., Li, H., Zhao, F., Chen, J., and Zhou, H. 2008.

Microstructure and mechanical properties of Mo/ DLCnanocomposite �lms. Diamond Relat. Mater. 17: 1949–1954.

362. Dai, M. J., Zhou, K. S., Lin, S. S. et al. 2007. Astudy on metal-doped diamond-like carbon �lm synthesized byion source and sputtering technique. Plasma Process. Polym.4: S215–S219.

363. Chen, X., Peng, Z., Fu, Z., Wu, S., Yue, W., and Wang,C. 2011. Microstructural, mechanical and tribologicalproperties of tungsten-gradually doped diamond-like carbon�lms with functionally graded interlayers. Surf. Coat.Technol. 205: 3631–3638.

364. Choy, K.-L. and Felix, E. 1999. Functionally gradeddiamond-like carbon coatings on metallic substrates.Mater. Sci. Eng. A 278: 162–169.

365. Monaghan, D. P., Teer, D. G., Logan, P. A., Efeoglu,I., and Arnell, R. D. 1993. Deposition of wear resistantcoatings based on diamond like carbon by unbalancedmagnetron sputtering. Surf. Coat. Technol. 60: 525–530.

366. Bertran, E., Corbella, C., Pinyol, A., Vives, M., andAndújar, J. L. 2003. Comparative study of metal/amorphous-carbon multilayer structures produced bymagnetron sputtering. Diamond. Relat. Mater. 12:1008–1012.

367. Kao, W. H. 2005. Tribological properties and highspeed drilling application of MoS 2 -Cr coatings. Wear258: 812–825. For catalogi g purp ses only

368. Chromik, R. R., Baker, C. C., Voevodin, A. A., andWahl, K. J. 2007. In situ tribometry of solid lubricantnanocomposite coatings. Wear 262: 1239–1252.

369. De Figueiredo, M. R., Muratore, C., Franz, R. et al.2010. In situ studies of TiC 1-x N x hard coatingtribology. Tribol. Lett. 40: 365–373.

370. Hu, J. J., Wheeler, R., Zabinski, J. S., Shade, P. A.,Shiveley, A., and Voevodin, A. A. 2008. Transmissionelectron microscopy analysis of Mo-W-S-Se �lm slidingcontact obtained by using focused ion beam microscope andin situ microtribometer. Tribol. Lett. 32: 49–57.

371. Dvorak, S. D., Wahl, K. J., and Singer, I. L. 2007. Insitu analysis of third body contributions to slidingfriction of a Pb-Mo-S coating in dry and humid air. Tribol.

Lett. 28: 263–274.

372. Scharf, T. W. and Singer, I. L. 2002. Role of thirdbodies in friction behavior of diamond-like nanocompositecoatings studied by in situ tribometry. Tribol. Trans. 45:363–371.

373. Singer, I. L., Dvorak, S. D., Wahl, K. J., and Scharf,T. W. 2003. Role of third bodies in friction and wear ofprotective coatings. J. Vac. Sci. Technol. A 21: S232–S240.

374. Muratore, C., Bultman, J. E., Aouadi, S. M., andVoevodin, A. A. 2011. In situ Raman spectroscopy forexamination of high temperature tribological processes.Wear 270: 140–145.

375. Scharf, T. W. and Singer, I. L. 2008. Third bodies andtribochemistry of DLC coatings. In Tribology ofDiamond-Like Carbon Films: Fundamentals and Applications,ed. C. Donnet and A. Erdemir, pp. 201–236. New York:Springer.

376. Scharf, T. W. and Singer, I. L. 2003. Monitoringtransfer �lms and friction instabilities with in situ Ramantribometry. Tribol. Lett. 14: 3–8.

377. Sawyer, W. G. and Wahl, K. J. 2008. Accessinginaccessible interfaces: In situ approaches to materialstribology. Mater. Res. Bull. 33: 1145–1150.

378. Wood, R. J. K. 2007. Tribo-corrosion of coatings: Areview. J. Phys. D: Appl. Phys. 40: 5502–5521.

379. Manhabosco, T. M. and Müller, I. L. 2009.Tribocorrosion of diamond-like carbon deposited on Ti6Al4V.Tribol. Lett. 33:193–197.

380. Landolt, D. 2006. Electrochemical and materialsaspects of tribocorrosion systems. J. Phys. D: Appl. Phys.39: 3121–3127.

381. Fusaro, R. L. 1994. Lubrication of space systems. NASATM 106392.

382. Roberts, E. W. 1990. Thin solid lubricant �lms inspace. Tribol. Int. 23: 95–104.

383. Todd, M. J. 1982. Solid lubrication of ball bearingsfor spacecraft mechanisms. Tribol. Int. 15: 331–337.

384. Christy, R. I. 1982. Dry lubrication for rollingelement spacecraft parts. Tribol. Int. 15: 265–271.

385. Stevens, K. T. and Todd, M. J. 1982. Parametric studyof solid-lubricant composites as ball-bearing cages.Tribol. Int. 15: 293–302.

386. Jones Jr, W. R. and Jansen, M. J. 2008. Tribology forspace applications. Proc. IMechE Part J: J. Eng. Tribol.222: 997–1004.

387. Fusaro, R. L. 2001. Preventing spacecraft failures dueto tribological problems. NASA TM-2001-210806.

388. McCook, N. L., Burris, D. L., Dickrell, P. L., andSawyer, W. G. 2005. Cryogenic friction behavior of PTFEbased solid lubricant composites. Tribol. Lett. 20:109–113.

389. Fleischauer, P. D. and Hilton, M. R. 1990.Applications of space tribology in the USA. Tribol. Int.23:135–139.

390. DellaCorte, C., Zaldana, A. R., and Radil, K. C. 2004.A systems approach to the solid lubrication of foil airbearings for oil-free turbomachinery. Tribol. Trans. 126:200–207.

391. DellaCorte, C. and Bruckner, R. J. 2010. Remainingtechnical challenges and future plans for oil-freeturbomachinery. NASA TM 2010-216762.

392. DellaCorte, C. and Edmonds, B. J. 1999.Self-lubricating composite containing chromium oxide.United States Patent No. 5,866,518.

393. DellaCorte, C., Lukaszewicz, V., Valco, M. J., Radil,K. C., and Heshmat, H. 2000. Performance and durabilityof high temperature foil air bearings for oil-freeturbomachinery. Tribol. Trans. 43: 774–780.

394. Wang, W. 2004. Application of a high temperatureself-lubricating composite coating on steam turbinecomponents. Surf. Coat. Technol. 177–178: 12–17.

395. DellaCorte, C. and Edmonds, B. J. 2009. NASA PS400: Anew high temperature solid lubricant coating for hightemperature wear applications. NASA TM-2009-215678.

396. Heshmat, H., Hryniewicz, P., Walton II, J. F., Willis,

J. P., Jahanmir, S., and DellaCorte, C. 2005. Lowfrictionwear-resistant coatings for high-temperature foil bearings.Tribol. Int. 38: 1059–1075.

397. Fanning, C. E. and Blanchet, T. A. 2008.High-temperature evaluation of solid lubricant coatings ina foil thrust bearing. Wear 265: 1076–1086.

398. Fox, V., Hampshire, J., and Teer, D. G. 1999. MoS 2/metal composite coatings deposited by closed-�eldunbalanced magnetron sputtering: Tribological propertiesand industrial uses. Surf. Coat. Technol. 112: 118–122.

399. Renevier, N. M., Fox, V. C., Teer, D. G., andHampshire, J. 2000. Performance of low friction MoS 2titanium composite coatings used in forming applications.Mater. Des. 21: 337–343.

400. Fox, V., Jones, A., Renevier, N. M., and Teer, D. G.2000. Hard lubricating coatings for cutting and formingtools and mechanical components. Surf. Coat. Technol. 125:347–353.

401. Renevier, N. M., Lobiondo, N., Fox, V. C., Teer, D.G., and Hampshire, J. 2000 Performance of MoS 2 / metalcomposite coatings used for dry machining and otherindustrial applications. Surf. Coat. Technol. 123: 84–91.

402. Amaro, R. I., Martins, R. C., Seabra, J. O., Renevier,N. M., and Teer, D. G. 2005. Molybdenum disulphide/titaniumlow friction coating for gears application. Tribol. Int.38: 423–434.

403. Wenlong, S., Jianxin, D., Hui, Z., and Pei, Y. 2010.Study on cutting forces and experiment of MoS 2 / Zr-coatedcemented carbide tool. Int. J. Adv. Manuf. Technol. 49:903–909.

404. Stoyanov, P., Chromik, R. R., Goldbaum, D., Lince, J.R., and Zhang, X. 2010. Microtribological performance ofAu-MoS 2 and Ti-MoS 2 coatings with varying contactpressure. Tribol. Lett. 40: 199–211.

405. Cho, M. H., Ju, J., Kim, S. J., and Jang, H. 2006.Tribological properties of solid lubricants (graphite, Sb2 S 3 , MoS 2 ) for automotive brake friction materials.Wear 260: 855–860.

406. Field, S. K., Jarratt, M., and Teer, D. G. 2004.Tribological properties of graphite-like and diamond-like

carbon coatings. Tribol. Int. 37: 949–956.

407. www.bekaert.com.

408. www.apnano.com.

409. Coldwell, H. L., Dewes, R. C., Aspinwall, D. K.,Renevier, N. M., and Teer, D. G. 2004. The use ofsoft/lubricating coatings when dry drilling BS L168aluminium alloy. Surf. Coat. Technol. 177−178: 716−726.

410. Wain, N., Thomas, N. R., Hickman, S., Wallbank, J.,and Teer, D. G. 2005. Performance of low-friction coatingsin the dry drilling of automotive Al-Si alloys. Surf. Coat.Technol. 200: 1885–1892.

411. Basnyat, P., Luster, B., Muratore, C. et al. 2008.Surface texturing for adaptive solid lubrication. Surf.Coat. Technol. 203: 73–79.

412. Voevodin, A. A. and Zabinski, J. S. 2006. Lasersurface texturing for adaptive solid lubrication. Wear 261:1285–1292.

413. Guleryuz, C. G. and Krzanowski, J. E. 2010. Mechanismsof self-lubrication in patterned TiN coatings containingsolid lubricant microreservoirs. Surf. Coat. Technol. 204:2392–2399.

414. Moshkovith, A., Per�liev, V., Gindin, D., Parkansky,N., Boxman, R., and Rapoport, L. 2007. Surface texturingusing pulsed air arc treatment. Wear 263: 1467–1469.

415. Zhang, X., Luster, B., Church, A. et al. 2009. Carbonnanotube-MoS 2 composites as solid lubricants. Appl.Mater. Int. 1: 735–739.

416. Dickrell, P. L., Sinnott, S. B., Hahn, D. W. et al.2005. Frictional anisotropy of oriented carbon nanotubesurfaces. Tribol. Lett. 18: 59–62.

417. Hu, J. J., Jo, S. H., Ren, Z. F., Voevodin, A. A., andJ. S. Zabinski. 2005. Tribological behavior andgraphitization of carbon nanotubes grown on 440C stainlesssteel. Tribol. Lett. 19: 119–125.

418. Joly-Pottuz, L. and Ohmae, N. Carbon-basednanolubricants. In Nanolubricants, ed. J.-M. Martin andN. Ohmae, pp. 93–147. England: John Wiley & Sons Ltd.

419. Street Jr., K. W., Miyoshi, K., and Vander Wal, R. L.2007. Application of carbon based nano-materials toaeronautics and space lubrication. In Superlubricity, ed.A. Erdemir and J.-M. Martin, pp. 311–340. Oxford:Elsevier. For cataloging purposes only

8 Chapter 8 - Processes andCharacterizations of Metal MatrixComposites

Andrewes, C. J. E., H. Y. Feng, and W. M. Lau. 2000.Machining of an aluminum/SiC composite using diamondinserts. Journal of Materials Processing Technology102:25–29.

Arsecularatne, J. A., L. C. Zhang, and C. Montross. 2006.Wear and tool life of tungsten carbide, PCBN and PCDcutting tools. International Journal of Machine Tools &Manufacture 46:482–491.

Arsenault, R. J., L. Wang, and C. R. Feng. 1991.Strengthening of composites due to microstructural changesin the matrix. Acta Metallurgica et Materialia 39:47–57.

Assar, A. E. M. 1999. Fabrication of metal matrix compositeby in�ltration process—Part 2: Experimental study. Journalof Materials Processing Technology 86:152–158.

Barmouz, M., P. Asadia, M. K. BesharatiGivi, and M.Taherishargh. 2010. Investigation of mechanical propertiesof Cu/SiC composite fabricated by FSP: Effect of SiCparticles’ size and volume fraction. Materials Science andEngineering A 528:1740–1749.

Blucher, J. T., U. Narusawa, M. Katsumata, and A. Nemeth.2001. Continuous manufacturing of �ber-reinforced metalmatrix composite wires—Technology and productcharacteristics. Composites Part A 32:1759–1766.

Brendel, A., V. Paffenholz, T. Köck, and H. Bolt. 2009.Mechanical properties of SiC long �bre reinforced copper.Journal of Nuclear Materials 386–388:837–840.

Caceres, C. H. and W. J. Poole. 2002. Hardness and �owstrength in particulate metal matrix composites. MaterialsScience and Engineering A 332:311–317.

Chambers, A. R. 1996. The machinability of light alloyMMCs. Composites: Part A 27:143–147.

Chandra, N. and H. Ghonem. 2001. Interfacial mechanics ofpush-out tests: Theory and experiments. Composites: Part A32:575–584.

Chang, S. T., W. H. Tuan, H. C. You, and I. C. Lin. 1999.Effect of surface grinding on the strength of NiAl and Al

2 O 3 /NiAl composites. Materials Chemistry and Physics59:220–224.

Chen, J., C. Hao, and J. Zhang. 2006. Fabrication of 3D-SiCnetwork reinforced aluminum–matrix composites bypressureless in�ltration. Materials Letters 60:2489–2492.

Choi, Y., M. E. Mullins, K. Wijayatilleke, and J. K. Lee.1992. Fabrication of metal matrix composites of TiC-Althrough self-propagating synthesis reaction. Metallurgicaland Materials Transactions 23A: 2387–2392.

Chou, Y. K., and J. Liu. 2005. CVD diamond tool performancein metal matrix composite machining. Surface & CoatingsTechnology 200:1872–1878.

Ciftci, I., M. Turker, and S. Ulvi. 2004. Evaluation oftool wear when machining SiCp reinforced Al-2014 alloymatrix composites. Materials and Design 25:251–255.

Davim, J. P. 2002. Diamond tool performance in machiningmetal–matrix composites. Journal of Materials ProcessingTechnology 128:100–105

Davim, J. P. and A. M. Baptista. 2000. Relationship betweencutting force and PCD cutting tool wear in machiningsilicon carbide reinforced aluminium. Journal of MaterialsProcessing Technology 103:417–423.

Degischer, H. P., P. Prader, and C. S. Marchi. 2001.Assessment of metal matrix composites forinnovations-intermediate report of a European thematicnetwork. Composites Part A: Applied Science andManufacturing 32(8):1161–1166.

D’Errico, G. E. and R. Calzavarini. 2001. Turning of metalmatrix composites. Journal of Materials ProcessingTechnology 119:257–260.

Diehl, W. and S. Detlev. 1990. Injection moulding ofsuperalloys and intermetallic phases. Metal Powder Report45:333–334, 336–338.

Ding, X., W. Y. H. Liew, and X. D. Liu. 2005. Evaluation ofmachining performance of MMC with PCBN and PCD tools. Wear259:1225–1234.

Doel, T. J. A. and P. Bowen. 1996. Effect of particle sizeand matrix aging condition on toughness of particlereinforced aluminium based metal matrix composites.

Materials Science and Technology 12:586–594.

El-Gallab, M. and M. Sklad. 1998a. Machining of Al:SiCparticulate metal matrix composites. Part I: Toolperformance. Journal of Materials Processing Technology83:151–158.

El-Gallab, M. and M. Sklad. 1998b. Machining of Al/SiCparticulate metal matrix composites Part II: Workpiecesurface integrity. Journal of Materials ProcessingTechnology 83:277–285.

El-Gallab, M. and M. Sklad. 2000. Machining of Al/SiCparticulate metal matrix composites Part III:Comprehensive tool wear models comprehensive tool wearmodels. Journal of Materials Processing Technology101:10–20.

Fei, W. D., M. Hu, and C. K. Yao. 2003. Effects of thermalresidual stress creeping on microstructure and tensileproperties of SiC whisker reinforced aluminum matrixcomposite. Materials Science and Engineering A 356:17–22.

Fiori, F., E. Girardin, A. Giuliani, T. Lorentzen, A.Pyzalla, F. Rustichelli, and V. Stanic. 2000. Neutrondiffraction measurements for the determination of residualstresses in MMC tensile and fatigue specimens. Physica B276–278:923–924.

Flom, Y. and R. J. Arsenault. 1989. Effect of particle sizeon fracture toughness of SiC/Al composite material. ActaMetall 37:2413–23.

García-Leiva, M. C., I. Ocaña, A. Martín-Meizoso, J. M.Martínez-Esnaola, V. Marqués, and F. Heredero. 2003. Hightemperature tensile and fatigue behaviour of theunidirectionally reinforced MMC Ti64/SiC. EngineeringFracture Mechanics 70:2137–2148.

Gariboldi, E., and C. Santulli, F. Stivali, and M. Vedani.1996. Evaluation of tensile damage in particulatereinforcedMMC’s by acoustic emission. Scripta Materialia 35:273–277.

German, R. M. and A. Bose. 1989. Fabrication ofintermetallic matrix composites. Materials Science andEngineering A 107:107–112.

Gooder, J. L. and S. Mall. 1995. Compressive behavior of atitanium metal matrix composite with a circular hole.Composite Structures 31:315–324.

Gotman, I., M. J. Koczak, and E. Shtessel. 1994.Fabrication of Al matrix in-situ composites viaself-propagating synthesis. Materials Science andEngineering A 187:189–199.

Hanabe, M. R. and P. B. Aswath. 1996. Al 2 O 3 /Alparticle-reinforced aluminum matrix composite bydisplacement reaction. Journal of Materials Research11:1562–1569.

Hanabe, M. and P. B. Aswath. 1997. Synthesis of in-situreinforced Al composites from Al-Si-Mg-O precursors. ActaMaterialia 45:4067–4076.

Harris, S. J., K. Dinsdale, Y. Gao, and B. Noble. 1988.Mechanical and physical behaviour of metallic and ceramiccomposites. In: Proc. 9th Risø Int. Sym. (eds. S.I.Andersen, H. Lilholt, O. B. Pederson) (Risø Press,Denmark): 373.

Hashim, J., L. Looney, and M. S. J. Hashmi. 2001. Theenhancement of wettability of SiC particles in castaluminium matrix composites. Journal of MaterialsProcessing Technology 119:329–335.

Heath, P. J. 2001. Developments in the applications of PCDtooling. Journal of Materials Processing Technology116:31–38.

Hong, S. Y. Y. Ding, and R. G. Ekkens. 1999. Improving lowcarbon steel chip breakability by cryogenic chip cooling.International Journal of Machine Tools & Manufacture39:1065–1085.

Hooper, R. M., J. L. Henshall, and A. Klopfer. 1999. Thewear of polycrystalline diamond tools used in the cuttingof metal matrix composites. International Journal ofRefractory Metals and Hard Materials 17:103–109.

Hu, M., W. D. Fei, and C. K. Yao. 2002. Effect of heattreatment on dislocation states and work hardeningbehaviors of SiCw/6061Al composite. Materials Letters56:637–641.

Huda, D., M. A. El-Baradie, and M. S. J. Hashmi. 1994.Analytical study for the stress analysis of metal matrixcomposites. Journal of Materials Processing Technology45:429–434.

Humphreys, F. J. 1991. The thermomechanical processing ofA1-SiC particulate composites. Materials Science andEngineering A 135:267–273.

Hung, N. P., F. Y. C. Boey, K. A. Khor, C. A. Oh, and H. F.Lee. 1995. Machinability of cast and powderformed aluminumalloys reinforced with SiC particles. Journal of MaterialsProcessing Technology 48:291–297.

Hung, N. P., F. Y. C. Boey, K. A. Khor, Y. S. Phua, and H.F. Lee. 1996. Machinability of aluminum alloys reinforcedwith silicon carbide particulates. Journal of MaterialsProcessing Technology 56:966–977.

Hung, N. P., N. L. Loh, and V. C. Venkatesh. 1999.Machining of metal matrix composites. In: Machiningof Ceramics and Composites. (eds. S. Jahanmir, M. Ramuluand P. Koshy) (Marcel Dekker, Inc., Basel, New York). Forcataloging purposes only

Hung, N. P., S. H. Yeo, K. K. Lee, and K. J. Ng. 1998. Chipformation in machining particle-reinforced metal matrixcomposites. Materials and Manufacturing Processes13(1):85–100.

Hung, N. P., Z. W. Zhong, and C. H. Zhong. 1997. Grindingof metal matrix composites reinforced with siliconcarbideparticles. Materials and Manufacturing Processes12(6):1075–1091.

Ilio, A. D., A. Paoletti, V. Tagliaferri, and F. Veniali.1996. An experimental study on grinding of silicon carbidereinforced aluminium alloys. International Journal ofMachine Tools and Manufacture 36:673–685.

Jawahir, I. S. 1990. On the controllability of chipbreaking cycles and modes of chip breaking in metalmachining. Annals of the CIRP 39(1):47–51.

Jiang, J. and B. Dodd. 1995. Workability of aluminium-basedmetal matrix composites in cold compression. Composites26:62–66.

Joshi, S. S., N. Ramakrishnan, and P. Ramakrishnan. 2001.Micro-structural analysis of chip formation duringorthogonal machining of Al/SiCp composites. Journal ofEngineering Materials and Technology 123:315–321.

Karthikeyan, R., G. Ganesan, R. S. Nagarazan, and B. C.Pai. 2001. A critical study on machining of Al/SiC

composites. Materials and Manufacturing Processes16(1):47–60.

Karakaş, M. S., A. Acir, M. Übeyli, and B. Ögel. 2006.Effect of cutting speed on tool performance in milling ofB4Cp reinforced aluminum metal matrix composites. Journalof Materials Processing Technology 178:241–246.

Kellie, J. and J. V. Wood. 1995. Reaction processing in themetals industry. Materials World 3:10–12.

Kilickap, E., O. Cakir, M. Aksoy, and A. Inan. 2005. Studyof tool wear and surface roughness in machining ofhomogenised SiC-p reinforced aluminium metal matrixcomposite. Journal of Materials Processing Technology164–165:862–867.

Kolednik, O. and K. Unterweger. 2008. The ductility ofmetal matrix composites- relation to local deformationbehavior and damage evolution. Engineering FractureMechanics 75:3663–3676.

Kuruvilla, A. K., K. S. Prasad, V. V. Bhanuprasad, and Y.R. Mahajan. 1990. Microstructure−property correlation in AlTiB 2 (XD) composites. Scripta Metallurgica et Materiala24:873–878.

Kwak, J. S. and Y. S. Kim, 2008. Mechanical properties andgrinding performance on aluminum-based metal matrixcomposites. Journal of Materials Processing Technology201:596–600.

Lee, R. S., G. A. Chen, and B. H. Hwang. 1995. Thermal andgrinding induced residual stresses in a silicon carbideparticle-reinforced aluminium metal matrix composite.Composites 26:425–429.

Lee, A.K., L.E. Sanchez-Caldera, S.T. Oktay, and N.P. Suh.1992. Liquid-metal mixing process tailors MMCmicrostructures. Advanced Materials & Processes 8:31–34.

Lee, H. S., J. S. Yeo, S. H. Hong, D. J. Yoon, and K. H.Na. 2001. The fabrication process and mechanical propertiesof SiCp/Al–Si metal matrix composites for automobileair-conditioner compressor pistons. Journal of MaterialsProcessing Technology 113:202–208.

Leggoe, J. W. 2004. Determination of the elastic modulus ofmicroscale ceramic particles via nanoindentation. Journalof Materials Research 19(8):2437–2447.

Leggoe, J. W., X. Z. Hu, M. V. Swain, and M. B. Bush. 1994.An ultra-micro indentation investigation of aspects of thefracture process in particulate reinforced metal matrixcomposites. Scripta Metallurgica et Materialia31(5):577–582.

Leng, J., G. Wu, Q. Zhou, Z. Dou, and X. L. Huang. 2008.Mechanical properties of SiC/Gr/Al composites fabricatedby squeeze casting technology. Scripta Materialia59:619–622.

Leucht, R. and H. J. Dudek. 1994. Properties of SiC �brereinforced titanium alloys processed by �bre coating andhot isostatic pressing. Materials Science and Engineering A188:201–210.

Lin, J. T., D. Bhattacharyya, and W. G. Ferguson. 1998.Chip formation in the machining of SiC particle reinforcedaluminium matrix composites. Composites Science andTechnology 58:285–291.

Lin, J. T., D. Bhattacharyya, and C. Lane. 1995. Case studymachinability of a silicon carbide reinforced aluminiummetal matrix composite. Wear 181–183:883–888.

Lin, C. B., Y. W. Hung, W. C. Liu, and S. W. Kang. 2001.Machining and �uidity of 356Al/SiC(p) composites. Journalof Materials Processing Technology 110:152–159.

Lin, Y., R.H. Zee, and B.A. Chin. 1991. In situ formationof three-dimensional TiC reinforcements in Ti-TiCcomposites. Metallurgical Transactions A 22:859–865.

Loh, N. H., S. B. Tor, and K. A. Khor. 2001. Production ofmetal matrix composite part by powder injection molding.Journal of Materials Processing Technology 108:398–407.

Ma, Z. Y., J. H. Li, S. X. Li, X. G. Ning, Y. X. Lu, and J.Bi. 1996. Property-microstructure correlation in-situformed Al 2 O 3 , TiB 2 and Al 3 Ti mixture-reinforcedaluminium composites. Journal of Materials Science 31(3):741–747.

Ma, Z. Y. and S.C. Tjong. 1997. In situ ceramicparticle-reinforced aluminum matrix composites fabricatedby reaction pressing in the TiO 2 (Ti)-Al-B (B 2 O 3 )systems. Metallurgical and Materials Transactions 28A(9):1931–1942.

Maity, P. C., S. C. Panigrahi, and P. N. Chakraborty. 1993.Preparation of aluminium-alumina in-situ particlecomposite by addition of titania to aluminium melt. ScriptaMetallurgica et Materialia 28:549–552.

McClintock, F. A. and A. S. Argon. 1966. MechanicalBehavior of Materials. (Addison-Wesley Publishing CompanyInc, Massachusetts, USA).

McDaniels, D. L. 1985. Analysis of stress–strain, fracture,and ductility behavior of aluminium matrix compositescontaining discontinuous silicon carbide reinforcements.Metallurgica Transactions 16A:1105–1115.

McLean, A., H. Soda, Q. Xia, A. K. Pramanick, A. Ohno, G.Motoyasu, T. Shimizu, S. A. Gedeon, and T. North. 1997.SiC particulate-reinforced, aluminium-matrix composite rodsand wires produced by a new continuous casting route.Composites Part A 28A:153–162.

Miracle, D. B. 2005. Metal matrix composites—From scienceto technological signi�cance. Composites Science andTechnology 65:2526–2540.

Mortensen, A. and J. A. Cornie. 1991. On the in�ltration ofmetal matrix composites. Metallurgical and MaterialsTransactions A 22:1160–1163.

Mussert, K. M., W. P. Vellinga, A. Bakker, and S. V. D.Zwaag. 2002. A nano-indentation study on the mechanicalbehaviour of the matrix material in an AA6061-Al 2 O 3MMC. Journal of Materials Science 37(4):789–794.

Ng, E. and D. K. Aspinwall. 2002. The effect of workpiecehardness and cutting speed on the machinability of AISIH13 hot work die steel when using PCBN tooling. Transactionof the ASME 124:588–594.

Oñoro, J., M. D. Salvadorb, and L. E. G. Cambroneroc. 2009.High-temperature mechanical properties of aluminium alloysreinforced with boron carbide particles. Materials Scienceand Engineering A 499:421–426.

Oxley, P. L. B. 1962. Shear angle solutions in orthogonalmachining. International Journal of Machine Tool Designand Research 2(3):219–229.

Ozben, T., E. Kilickap, and O. Çakır. 2008. Investigationof mechanical and machinability properties of SiC particlereinforced Al-MMC. Journal of Materials Processing

Technology 198:220–225.

Ozcatalbas, Y. 2003a. Investigation of the machinabilitybehaviour of Al4C3 reinforced Al-based composite producedby mechanical alloying technique. Composites Science andTechnology 63(1):53–61.

Ozcatalbas, Y. 2003b. Chip and built-up edge formation inthe machining of in situ Al 4 C 3 –Al composite. Materials& Design 24(3):215–221.

Pandey, A. B., B. S. Majumdar, and D. B. Miracle. 2000.Deformation and fracture of a particle-reinforced aluminumalloy composites: Part I. Experiments. MetallurgicalTransactions 31A:921–936.

Partridge, P. G. and C. M. Ward-Close. 1993. Processing ofadvanced continuous �bre composites: Current practice andpotential developments. International Materials Reviews38:1–24.

Pedersen, W. and M. Ramulu. 2006. Facing SiCp/Mg metalmatrix composites with carbide tools. Journal of MaterialsProcessing Technology 172:417–423.

Pelleg, J., S. Sofer, M. Alkoby, and M. Ganor. 1995.In-situ formation of the Ti 3 A1-SiC metal matrix compositesystem from its components. Materials Science andEngineering A 203:305–313.

Peng, J., D. H. W. Li, J. D. Y. Xie, and G. Liu. 2008.Study on the yield behavior of Al 2 O 3 –SiO 2 (sf)/Al–Simetal matrix composites. Materials Science and EngineeringA 486:427–432.

Pereyra, R. and Y. L. Shen. 2004. Characterization ofparticle concentration in indentation-deformed metalceramiccomposites. Material Characterization 53:373–380.

Pereyra, R. and Y. L. Shen. 2005. Characterization ofindentation-induced “particle crowding” in metal matrixcomposites. International Journal of Damage Mechanics14(3):197–213.

Pramanik, A. and L. C. Zhang. 2009. Machining metal matrixcomposites (Chapter 5). In: Machining Composite Materials.(ed. J. P. Davim) (John Wiley & Sons, United Kingdom)(ISBN: 9781848211704).

Pramanik, A., L. C. Zhang, and J. A. Arsecularatne. 2007.

An FEM investigation into the behaviour of metal matrixcomposites: Tool−particle interaction during orthogonalcutting. International Journal of Machine Tools andManufacture 47:1497–1506.

Pramanik, A., L. C. Zhang, and J. A. Arsecularatne. 2008a.Deformation mechanisms of MMCs under indentation.Composites Science and Technology 68:1304–1312.

Pramanik, A., L. C. Zhang, and J. A. Arsecularatne. 2008b.Machining of metal matrix composites: Effect of ceramicparticles on residual stress, surface roughness and chipformation. International Journal of Machine Tools &Manufacture 48:1613–1625.

Prasad, V. V. B., B. V. R. Bhat, Y. R. Mahajan, and P.Ramakrishnan. 2002. Structure-property correlation indiscontinuously reinforced aluminium matrix composites as afunction of relative particle size ratio. MaterialsScience and Engineering A 337:179–186. For cataloging purps s only

Quan, Y. and Z. Zehua. 2000. Tool wear and its mechanismfor cutting SiC particle reinforced aluminium matrixcomposites. Journal of Materials Processing Technology100(1):194–199.

Quan, Y. M., Z. H. Zhou, and B. Y. Ye. 1999. Cuttingprocess and chip appearance of aluminum matrix compositesreinforced by SiC particle. Journal of Materials ProcessingTechnology 91(1):231–235.

Quan, Y. and B. Ye. 2003. The effect of machining on thesurface properties of SiC/Al composites. Journal ofMaterials Processing Technology 138(1–3):464–467.

Rashad, R. M. and T. M. El-Hossainy. 2006. Machinability of7116 structural aluminum alloy. Materials andManufacturing Processes 21:23–27.

Ralph, B., H. C. Yuen, and W. B. Lee. 1997. The processingof metal matrix composites—An overview. Journal ofMaterials Processing Technology 63:339–353.

Ranganath, S., M. Vijayakumar, and J. Subrahmanyam. 1992.Combustion-assisted synthesis of TiTiBTiC composite viathe casting route. Materials Science and Engineering A149:253–257.

Reddy, N. S. K., S. Kwang-sup, and M. Yang. 2008.

Experimental study of surface integrity during end millingof Al/SiC particulate metal–matrix composites. Journal ofMaterials Processing Technology 201:574–579.

Robertson, D. D. and S. M. Mall. 1994. Micromechanicalanalysis for thermoviscoplastic behavior of unidirectional�brous composites. Composites Science and Technology50:483–496.

Ronald, B. A., L. Vijayaraghavan, and R. Krishnamurthy2009. Studies on the in�uence of grinding wheel bondmaterial on the grindability of metal matrix composites.Materials and Design 30:679–686.

Sahoo, P. and M. J. Koczak. 1991a. Analysis of in situformation of titanium carbide in aluminum alloys.Materials Science and Engineering A 144:37–44.

Sahoo, P. and M. J. Koczak. 1991b. Microstructure−propertyrelationships of in situ reacted TiC/AlCu metal matrixcomposites. Materials Science and Engineering A 131:69–76.

Sharma, S. R., Z. Y. Ma, and R.S. Mishra. 2004. Effect offriction stir processing on fatigue behavior of A356alloy. Scripta Materialia 51:237–241.

Shen, Y. L. and N. Chawla. 2001b. On the correlationbetween hardness and tensile strength in particlereinforced metal matrix composites. Materials Science andEngineering A 297(1–2):44–47.

Shen, Y. L., E. Fishencord, and N. Chawla. 2000.Correlating macrohardness and tensile behavior indiscontinuously reinforced metal matrix composites. ScriptaMaterialia 42(5):427–432.

Shen, Y. L., J. J. Williams, G. Piotrowski, N. Chawla, andY. L. Guo. 2001a. Correlation between tensile andindentation behavior of particle-reinforced metal matrixcomposites: An experimental and numerical study. ActaMaterialia 49:321–3229.

Shi, N. and R. J. Arsenault. 1991. Analytical evaluation ofthe thermal residual stresses in Si/Al composites. JSMEInternational Journal 34(2):143–155.

Slipenyuk, A., V. Kuprin, Y. Milman, J. E. Spowart, and D.B. Miracle. 2004. The effect of matrix to reinforcementparticle size ratio (PSR) on the microstructure andmechanical properties of a P/M processed AlCuMn/SiCp MMC.

Materials Science and Engineering A 381:165–170.

Songmene, V. and M. Balazinsk. 1999. Machinability ofgraphitic metal matrix composites as a function ofreinforcing particles. Annals of the CIRP 48(1):77–80.

Sreejith, P. S. 2006. Tool wear of binderless PCBN toolduring machining of particulate reinforced MMC. TribologyLetters 22(3):265–269.

Stephens J. J., J. P. Lucas, and F. M. Hosking. 1988. CastAl–7Si composites: effect of particle type and size onmechanical properties. Scripta Materialia 22:1307–1312.

Subramanian, P. R., S. Krishnamurthy, S. T. Keller, and M.G. Mendiratta. 1998. Processing of continuously reinforcedTi-alloy metal matrix composites (MMC) by magnetronsputtering. Materials Science and Engineering A 244:1–10.

Syn, C. K., D. R. Lesuer, and O. D. Sherby. 1996. Enhancingtensile ductility of a particulate-reinforced aluminummetal matrix composite by lamination with Mg-9%Li alloy.Materials Science and Engineering A 206:201–207.

Taheri-Nassaj, E., M. Kobashi, and T. Choh. 1996.Fabrication and analysis of an in situ TiB2/Al composite byreactive spontaneous in�ltration. Scripta Materialia34:1257–1265.

Taheri-Nassaj, E., M. Kobashi, and T. Choh. 1997.Fabrication and analysis of in situ formed boride/AIcomposites by reactive spontaneous in�ltration. ScriptaMaterialia 37:605–614.

Thomas, M. P. and J. E. King. 1996. Improvement of themechanical properties of 2124 Al/SiC, MMC plate byoptimisation of the solution treatment. Composites Scienceand Technology 56:1141–1149.

Thomas, M. P. and M. R. Winstone. 1999. Longitudinalyielding behaviour of SiC-�bre-reinforced titaniummatrixcomposites. Composites Science and Technology 59:297–303.

Tjong, S. C. and Z. Y. Ma. 2000. Microstructural andmechanical characteristics of in situ metal matrixcomposites. Materials Science and Engineering 29:49–113.

Übeyli, M., A. Acir, M. S. Karakaş, and B. Ögel. 2007.Study on performance of uncoated and coated tools inmilling of Al-4%Cu/B4C metal matrix composites. Materials

Science and Technology 23(8):945–950.

Übeyli, M., A. Acir, M. S. Karakaş, and B. Ögel. 2008.Effect of feed rate on tool wear in milling of Al-4%Cu/B4Cp composite. Materials and Manufacturing Processes23:865–870.

Vassel, A. 1999. Continuous �bre reinforced titanium andaluminium composites: A comparison. Materials Science andEngineering A 263:305–313.

Varadarajana, Y. S., L. Vijayaraghavan, and R.Krishnamurthy. 2006. Performance enhancement throughmicrowave irradiation of K20 carbide tool machining Al/SiCmetal matrix composite. Journal of Materials ProcessingTechnology 173(2):185–193.

Wang, Z., T. Chen, and D. J. Lloyd. 1993. Stressdistribution in particulate-reinforced metal-matrixcomposites subjected to external load. MetallurgicalTransactions 24(A):197–207.

Ward-Close, C. M. and P. G. Partridge. 1990. A �bre coatingprocess for advanced metal-matrix composites. Journal ofMaterials Science 25:4315–4323.

Warren, R. 1988. Mechanical and physical behaviour ofmetallic and ceramic composites. In: Proc. 9th RisøInt.Sym. (eds. S. I. Andersen, H. Lilholt, O. B. Pederson)(Risø Press, Denmark): 233.

Weinert, K. 1993. A consideration of tool wear mechanismwhen machining metal matrix composite (MMC). Annals ofCIRP 42:95–98.

Xiandong, S., L. Chengping, L. Zhuoxuan, and O. Liuzhang.1997. The Fabrication and properties of particlereinforced cast metal matrix composites. Journal ofMaterials Processing Technology 63:426–431.

Xu, H. and E. J. Palmiere. 1999. Particulate re�nement andredistribution during the axisymmetric compression of anAl/SiCp metal matrix composite. Composites: Part A30:203–211.

Yea, H., X. Y. Liu, and H. Hong. 2008. Fabrication of metalmatrix composites by metal injection molding—A review.Journal of Materials Processing Technology 200:12–24.

Yih, P. and D. D. L. Chung. 1997. A comparative study of

the coated �ller method and the admixture method of powdermetallurgy for making metal–matrix composites. Journal ofMaterials Science 32:2873–2882.

Yong, M. S. and A. J. Clegg. 2005. Process optimisation fora squeeze cast magnesium alloy metal matrix composite.Journal of Materials Processing Technology 168:262–269.

Yuen, H. C., W. B. Lee, and B. Ralph. 1995. Hot-rollbonding of aluminium matrix composites with differentvolume fractions of alumina. Journal of Materials Science30:843– 848.

Zhang, F., P. Sun, and X. Li. 2001. A comparative study onmicroplastic deformation behavior in a SiCp/2024Alcomposite and its unreinforced matrix alloy. MaterialsLetters 49:69–74.

Zhao, B., C. S. Liu, X. S. Zhu, and K. W. Xu. 2002.Research on the vibration cutting performance of particlereinforced metallic matrix composites SiCp/Al. Journal ofMaterials Processing Technology 129(1–3):380–384.

Zhong, Z. W. 2003. Grinding of aluminium-based metal matrixcomposites reinforced with Al 2 O 3 or SiC particles.International Journal of Advanced Manufacturing Technology21:79–83.

Zhong, Z. and N. P. Hung. 2000. Diamond turning andgrinding of aluminium based metal matrix composites.Materials and Manufacturing Processes 15:853–865.

Zhong, Z. and N. P. Hung. 2002. Grinding ofalumina/aluminum composites. Journal of MaterialsProcessing Technology 123:13–17. For cataloging purpos sonly

9 Chapter 9 - Processing Science forPolymeric Composites in Aerospace

Abdalla, FH, and Mutasher, SA 2007, Design and fabricationof low cost �lament winding machine, Materials and Design,28(1), 234–239.

Advani, SG 1994, Resin transfer molding �ow phenomena inpolymeric composites, in In Flow and Rheology in PolymerComposite Manufacturing, Elsevier Science, Newark, pp.465–515.

Ahn, SH, Lee, WI, and Springer, GS 1995, Measurement of thethree-dimensional permeability of �bre preforms usingembedded �bre optic sensors, Journal of CompositeMaterials, 29(6), 714–733.

Anamateros, E, Leone, A, and Severoni, R 2001, Manufactureof an Helicopter Structure with Resin Transfer molding,presented at the Specialists Meeting on Low Cost CompositeStructures, Loen, Norway.

ANSYS revision 5.0, Theory & User’s Manuals, SwansonAnalysis Systems Inc., Houston, USA.

Baden Fuller, AJ 1973, Engineering Field Theory, 11th edn,Pergamon Press, Oxford, New York.

Batch, GL, and Macosko, CW 1993, Heat transfer and cure inpultrusion: Model and experimental veri�cation, AmericanInstitute of Chemical Engineers Journal, 39(7), 1228–1241.

Blest, DC, and McKee, S 1999, Curing simulation byautoclave resin infusion, Composites Science andTechnology, 59(16), 2297–2313.

Bogetti, TA, and Gillespie Jr, JW 1991, Two-dimensionalcure simulation of thick thermosetting composites, Journalof Composite Materials, 25(3), 239–273.

Borchardt, JK 2004, Unmanned aerial vehicles spurcomposites use, Journal of Reinforced Plastics, 48(4),28–31.

Brouwer, WD, Van Herpt, ECFC, and Labordus, M 2003, Vacuuminjection moulding for large structural applications,Composites Part A: Applied Science and Manufacturing,34(6), 551–558.

Calhoun, DR, Yalvac, S, Wetters, DG, and Raeck, CA 1996,

Critical issues in model veri�cation for the resintransfer molding process, Polymer Composites, 17(1), 11–22.

Campbell, FC, Mallow, AR, and Browning, CED 1995, Porosityin carbon �bre composites: An overview of causes, Journalof Advanced Materials, 26(4), 18–32.

Chachad, YR, Roux, JA, and Vaughan, JG 1996, Thermal modelfor three-dimensional irregular shaped pultruded �berglasscomposites, Composite Materials, 30(6), 692–721.

Choi, JH, and Lee, DG 1995, Expert cure system for thecarbon �bre epoxy composite materials, Journal ofComposite Materials, 29(9), 1181–1201.

Dasgupta, A, and Agarwal, RK 1992, Orthotropic thermalconductivity of plain weave fabric composites, Journal ofComposite Materials, 26(18), 2736–2758.

Dusi, MR, Lee, WI, Ciriscioli, PR, and Springer, GS 1987,Cure kinetics and viscosity of Fiberite 976 resin, Journalof Composite Materials, 21(3), 243–261.

Garg, A 1993, Resin Transfer Moulding of AerospaceComposite Structures—A Review, CRC-AS TR93003, CRC-AS,Australia.

Groppe, D 2003, Robots speed lay-up of composite materials,Journal of Reinforced Plastics, 47(4), 44–47.

Grove, SM 1988, Thermal modelling of tape laying withcontinuous carbon �bre-reinforced thermoplastic,Composites, 19(5), 367–375.

Hackett, RM, and Zhu, SZ 1992, Two-dimensional �niteelement model of the pultrusion process, Journal ofReinforced Plastics and Composites, 11(12), 1322–1351.

Halley, P, and Mackey, ME 1996, Chemorheology ofthermosets—An overview, Polymer Engineering and Science,36(5), 593–609.

Hammami, A 2002, Key factors affecting permeabilitymeasurement in the vacuum infusion molding process,Polymer Composites, 23(6), 1057–1067.

Heitzmann, KF 1992, Determination of in Plane Permeabilityof Woven and Non-woven Fabrics, Master’s thesis,University of Illinois, Urbana, Chicago.

Johnson, RJ, and Pitchumani, R 2003, Enhancement of �ow inVARTM using localized induction heating, CompositesScience and Technology, 63(15), 2201–2215.

Jones, RM 1975, Mechanics of Composite Materials,Mc-Graw-Hill Kogakusha Ltd, Tokyo.

Joseph, B, Hanratty, FW, and Kardos, JL 1995, Model basedcontrols of voids and product thickness during autoclavecuring of Carbon/Epoxy composite laminates, Journal ofComposite Materials, 29(8), 1000–1023.

Joshi, SC, and Chen, X 2011, Time-variant simulation ofmulti-material thermal pultrusion, Applied CompositeMaterials, 18(4), 283–296.

Joshi, SC, and Lam, YC 2001, Three-dimensional FE/NCVsimulation of pultrusion process with temperaturedependentmaterial properties including resin shrinkage, CompositesScience and Technology, 61(11), 1539–1547.

Joshi, SC 1998, Investigations on Prepreg Moulding, RTM,and RFI Processes of Composites Manufacturing, PhD thesis,Monash University, Clayton campus, Australia.

Joshi, SC 2009, Pragmatism in semi-steady modular�nite-grid simulation methodology for aerospace composites,Simulation Modelling Practice and Theory (Elsevier), 17(5),839–849.

Joshi, SC, Lam YC, and Liu, XL 2000, Mass conservation innumerical simulation of resin �ow, Composites Part A:Applied Science and Manufacturing, 31(10), 1061–1068.

Joshi, SC, Liu, XL, and Lam, YC 1999, A numerical approachfor modelling of polymer curing in �bre reinforcedcomposites, Composites Science and Technology, 59(7),1003–1013.

Joshi, SC, Liu, XL, Lam, YC, and Sheridan, J 1999,Simulation of resin �lm infusion process using �niteelement/nodal control volume approach, Advanced CompositesLetters, 8(3), 101–104.

Kang, MK, Lee, WI, Yoo, JY, and Cho, SM 1995, Simulation ofmold �lling process during resin transfer molding,Materials Processing and Manufacturing science, 3, 297–313.

Kinsey, SP, Sheikh, AH, and Lou, YS 1992, A thermal modelfor cure of advanced composite materials, First

International Conference on Transport Phenomena inProcessing, Lancaster, pp. 1353–1368.

Kranbuehl, DE 1995, FDEMS Sensing for Automated IntelligentProcessing of Polyimides, ASME Materials Division, Vol.69, pp. 1017–1030.

Krishnaswamy, P 1994, Durability of Composite Tools, CRC-ASTR 94007, CRC-ACS, Melbourne, Australia.

Lam, YC, Joshi, SC, and Liu, XL 1998, Numerical simulationof mould-�lling process in resin transfer moulding,Composites Science and Technology, 60(6), 845–855.

Lee, SM 1991, Tooling, in International Encyclopedia ofComposites, Vol. 6, VCH Publishers, New York, pp. 483–489.

Lee, WI, and Springer, GS 1987, A model of themanufacturing process of thermoplastic matrix composites,Journal of Composite Materials, 21(11), 1017–1055.

Lee, WI, Loos, AC, and Springer, GS 1982, Heat of reaction,degree of cure, and viscosity of Hercules 3501-6 resin,Journal of Composite Materials, 16(6), 510–520.

Leonard-Williams, S 2008, The crossover from RTM to resininfusion, Reinforced Plastics, 52(10), 28–29.

Letterman, LE 1986, Resin Film Infusion Process andApparatus, Patent No. 4,622,091, The Boeing Company,Seattle, Washington.

Li, JH, Joshi, SC, and Lam, YC 2002, Curing optimizationfor pultruded composite sections, Composites Science andTechnology, 62(3), 457–467.

Lin, M, and Hahn, HT 1996, A robust and ef�cient approachfor RTM Simulation, in Abstracts on Advanced Materials:Development, Characterization, Processing, and MechanicalBehavior, Vol. 74, Materials Division, American Societyfor Mechanical Engineers, USA. For cataloging purposesonly

Loos, AC, and MacRae, JD 1996, A process simulation modelfor the manufacture of a blade-stiffened panel by theresin �lm infusion process, Composites Science andTechnology, 56(3), 273–289.

Loos, AC, and Springer, GS 1983, Curing of epoxy matrixcomposites, Journal of Composite Materials, 17(2),

135–169.

Lossie, M, and Van Brussel, H 1994, Design principles in�lament winding, Composites Manufacturing, 5(1), 5–13.

LUSAS (version 13.0), Theory & User’s manuals, FEA Limited,Surrey, UK.

MacRae, JD, Loos, AC, Deaton, JW, and Hasko, GH 1994,Development and Veri�cation of a Resin Film Infusion/ResinTransfer Molding Simulation Model For Fabrication ofAdvanced Textile Composites, NASA-CR-197439, Virginiapolytechnic institute and state university, USA.

Maier, RS, Rohaly, TF, Advani, SG, and Fickie, KD 1996, Afast numerical method for isothermal resin transfer mold�lling’, International Journal of Numerical Methods andEngineering, 39, 1405–1417.

Mallick, PK 1993, Fibre-reinforced Composites: Materials,Manufacturing, and Design, 3rd edn, Marcel Dekker, Inc.,New York.

Marsh, G 1997, Putting SCRIMP in context, ReinforcedPlastics, 41(1), 22–26.

Niu, MCY 1992, Tooling, in Composite Air Frames—PracticalDesign Information And Data, Conmilit Press Ltd, HongKong, pp. 127–175.

Ozawa, T 1971, Kinetics of non-isothermal crystallization,Polymer, 12(3), 150–158.

Payne, JB 1988, Determining cure cycles for thermosettingepoxy resins, Fabricating Composites 1988, Society ofManufacturing Engineers, Philadelphia, Pennsylvania, pp.1–10.

Perry, MJ 1993, Analysis of Resin Transfer Molding:Material Characterization, Molding and Simulation, PhDthesis, Ohio State University, USA.

Rice, EV, Owen, MJ, and Pickering, SJ 1994, Simultaneousengineering of components and tooling for resin transfermoulding (RTM) composites, Engineering Systems design andanalysis, 64(2), 269–283.

Ruan, Y, and Liu, J 1994, A steady state heat transfermodel for �ber-reinforced-thermoplastic pultrusion processusing the �nite element method, Journal of Materials

Processing & Manufacturing Science, 3(1), 91–113.

Rudd, CD, Long, AC, Kendall, KN, and Mangin, C 1997, LiquidMoulding Technologies, 1st edn, Woodhead publishinglimited, Cambridge, England.

Sharma, D, McCarty, TA, Roux, JA, and Vaughan, JG 1998,Investigation of dynamic pressure behaviour in apultrusion die, Journal of Composite Materials, 32(1),929–950.

Shim, SB, Ahn, K, and Seferis, JC 1994, Flow and voidcharacterisation of stitched structural composites usingresin �lm infusion process (RFIP), Polymer Composites,15(6), 453–463.

Shim, SB, Ahn, K, Seferis, JC, Berg, AJ, and Hudson, W1995, Cracks and microcracks in stitched structuralcomposites manufactured with resin �lm infusion process,Journal of Advanced Materials, 26(4), 48–62.

Sumerak, JE 1997, The pultrusion process for automatedmanufacture of engineered composite pro�les, in Mallick,PK, Composites Engineering, Chapter 11, pp. 549–577.

Suratno, BR, Ye, L, and Mai, YW 1998, Simulation oftemperature and curing pro�les in pultruded compositerods, Composite Science and Technology, 58(2), 191–197.

Twardowski, TE, Lin, SE, and Geil, PH 1993, Curing in thickcomposite laminates: Experiment and simulation, Journal ofComposite Materials, 27, 216–250.

Vaidya, UK, Abraham, A, and Bhide, S 2001, Affordableprocessing of thick section and integral multi-functionalcomposites, Composites Part A: Applied Science andManufacturing, 32(8), 1133–1142.

Valliappan, M, Roux, JA, Vaughan, JG, and Arafat, ES 1996,Die and post-die temperature and cure in Graphite/ Epoxycomposites, Composites: Part B, 27(1), 1–9.

Varma, RR, and Advani, SG 1994, Three-dimensionalsimulations of �lling in resin transfer molding, Advancesin �nite element analysis in ©uid dynamics (ASME), 200,21–27.

Vergnaud, JM, and Bouzon, J 1992, Numerical analysis for aplane sheet: One dimensional heat transfer and curereaction, in In Cure of Thermosetting Resins: Modelling and

Experiments, Springer-Verlag, London, pp. 105–107.

Viola, G, Portwood, T, Ubrich, P, and DeGroot, HR 1990,Numerical optimization of pultrusion line operatingparameters, 35th International SAMPE Symposium, CA, USA,pp. 1968–1975.

Vodicka, R, and Evans, N 1994, User’s Manual for Cure aComputer Simulation of the Autoclave Curing Process forComposite Layup, CRC-AS-PS-121–001, CRC-AS, Melbourne,Australia.

Weitzezenböck, JR, Shenoi, RA, and Wilson, PA 1997,Measurement of three-dimensional permeability, CompositesPart A, 29(1–2), 159–169.

Wetherhold, RC, and Wang, JZ 1994, Predicting transversethermal conductivity of continuous �bre composites, Journalof Composite Materials, 28(15), 1491–1498.

Wilcox, JAD, and Wright, DT 1998, Towards pultrusionprocess optimization using arti�cial neural networks,Journal of Materials Processing Technology, 83(1–3),131–141.

Young, WB 1994, Three-dimensional non-isothermal mold�lling simulations in resin transfer molding, PolymerComposites, 15(2), 118–127.

Young, WB 1995, Thermal behavior of the resin and mold inthe process of resin transfer molding, Reinforced Plasticsand Composites, 14(4), 310–331.

Young, WB 2009, Development of a helicopter landing gearprototype using resin infusion molding, Journal ofReinforced Plastics and Composites, 28(7), 833–849. Forcataloging purposes only

10 Chapter 10 - CarbonNanotube-Reinforced Polymer Compositesfor Aerospace Application

1. Jeevananda, T., Jang, Y.K., Lee, J.H., Siddaramaiah, D.,Urs M.V., and Ranganathaiah, C. 2009. Investigation ofmulti-walled carbon nanotube reinforced high-densitypolyethylene/carbon black nanocomposites using electricalDSC and positron lifetime spectroscopy techniques. PolymInt 58:755–80.

2. Li, Q., Siddaramaiah, Kim, N.H., Yoo, G.H., and Lee,J.H. 2009. Positive temperature coef�cient characteristicand structure of graphite nano�bers reinforced high-densitypolyethylene/carbon black nanocomposites. Compos Part B40:218–24.

3. Cheng, H.K.F., Sahoo, N.G., Khin, T.H., Li, L., Chan,S.H., Zhao, J., and Chen, G. 2010. The role offunctionalized carbon nanotubes in a PA6/LCP blend. JNanosci Nanotech 10:5242–51.

4. Sahoo, N.G., Cheng, H.K.F., Li, L., Chan, S.H., Judeh,Z., and Zhao, J. 2009. Speci�c functionalization of carbonnanotubes for advanced polymer nanocomposites. Adv FunctMater 19:3962–71.

5. Sahoo, N.G., Cheng, H.K.F., Cai, J., Li., L., Chan,S.H., Zhao, J., and Yu, S. 2009. Improvement of mechanicaland thermal properties of carbon nanotube compositesthrough nanotube functionalization and processing methods.Mater Chem Phys 117:313–20.

6. Sahoo, N.G., Cheng, H.K.F., Pan, Y., Li, L., Chan, S.H.,and Zhao, J. 2010. Strengthening of liquid crystallinepolymer by functionalized carbon nanotubes throughinterfacial interaction and homogeneous dispersion. PolymAdv Tech, DOI: 10.1002/pat.1704.

7. Wang, H., Zhang, H., Zhao, W., Zhang, W., and Chen, G.2008. Preparation of polymer/oriented graphite nanosheetcomposite by electric �eldinducement. Compos Sci Technol68:238–43.

8. Yu, A., Ramesh, P., Itkis, M.E., Elena, B., and Haddon,R.C. 2007. Graphite nanoplatelet-epoxy composite thermalinterface materials. J Phys Chem C 111:7565–9.

9. Debelak, B. and Lafdi, K. 2007. Use of exfoliatedgraphite �ller to enhance polymer physical properties.

Carbon 45:1727–34.

10. Chen, X., Zheng, Y.P., Kang, F., and Shen, W.C. 2006.Preparation and structure analysis of carbon/ carboncomposite made from phenolic resin impregnation intoexfoliated graphite. J Phys Chem Solids 67:1141–4.

11. Khanna, V. and Bakshi, B.R. 2009. Carbon nano�berpolymer composites: Evaluation of life cycle energy use.Environ Sci Technol 43:2078–84.

12. Tibbetts, G.G, Lake, M.L., Strong, K.L., and Rice, B.P.2007. A review of the fabrication and properties ofvapor-grown carbon nano�ber/polymer composites. Compos SciTechnol 67:1709–18.

13. Chipara, M., Lozano, K., Hernandez, A., and Chipara, M.2008. TGA analysis of polypropylene-carbon nano�berscomposites. Polym Degrad Stab 93:871–6.

14. Sahoo, N.G., Rana, S., Cho, J.W., Li, L., and Chan,S.H. 2010. Polymer nanocomposites based on functionalizedcarbon nanotubes. Prog Polym Sci 35:837–67.

15. Xie, X. L., Mai, Y.W., and Ping, X. 2005. Dispersionand alignment of carbon nanotubes in polymer matrix: Areview. Mater. Sci. Eng.: R: Reports 49:89–112.

16. Spitalsky, Z., Dimitrios, T., Papagelis, K., andGaliotis, C. 2010. Carbon nanotube-polymer composites:Chemistry, processing, mechanical and electricalproperties. Prog Polym Sci 35:357.

17. Saito, R., Dresselhaus, D., and Dresselhaus, M.S. 1999.Physical Properties of Carbon Nanotubes. London: ImperialCollege Press.

18. Dresselhaus, M.S., Lin, Y.M., Rabin, O. et al. 2003.Nanowires and nanotubes. Mater Sci Eng C 23:129–40.

19. Iijima, S. and Ichihashi, T. 1993. Single-shell carbonnanotubes of 1-nm diameter. Nature 363:603–5.

20. Guo, T., Nikolaev, P., Thess, A., Colbert, D.T., andSmalley, R.E. 1995. Catalytic growth of single-wallednanotubes by laser vaporization. Chem Phys Lett 243:49–54.

21. Nikolaev, P., Bronikowski, M.J., Bradley, R.K. et al.1999. Gas-phase catalytic growth of single-walled carbonnanotubes from carbon monoxide. Chem Phys Lett 313:91–7.

22. Joseyacaman, M., Mikiyoshida,M., Rendon, L., andSantiesteban, J.G. 1993. Catalytic growth of carbonmicrotubules with fullerene structure. Appl Phys Lett62:657–9.

23. Huang, S.M., Woodson, M., Smalley, R.E., and Liu, J.2004.Growth mechanism of oriented long single walledcarbon nanotubes using fast-heating chemical vapordeposition process. Nano Lett 4:1025–8.

24. Chae, H.G., Liu, J., and Kumar, S. 2006. Carbonnanotubes properties and applications. In: O’Connell MJ,editor. Carbon Nanotube-Enabled Materials. Boca Raton:Taylor & Francis Group, LLC. For ataloging purp ses only

25. Ma, C., Zhang, W., Zhu, Y. et al. 2008. Alignment anddispersion of functionalized carbon nanotubes in polymercomposites induced by an electric �eld. Carbon 46:706–10.

26. Ajayan, P.M., Schadler, L.S., and Braun, P.V. 2003.Nanocomposite Science and Technology. Weinheim, Germany:Wiley-VCH, GmbH & Co. KgaA.

27. Lu, J.P. 1997. Elastic properties of carbon nanotubesand nanoropes. Phys Rev Lett 79:1297–300.

28. Yu, M.F., Lourie, O., Dyer, M.J., Moloni, K., Kelly,T.F., and Ruoff, R.S. 2000. Strength and breakingmechanism of multiwalled carbon nanotubes under tensileload. Science 287:637–40.

29. Ajayan, P.M., Stephan, O., Colliex, C., and Trauth, D.1994. Aligned carbon nanotube arrays formed by cutting apolymer resin-nanotube composite. Science 265:1212–4.

30. Paredes, J.I. and Burghard, M. 2004. Dispersions ofindividual single-walled carbon nanotubes of high length.Langmuir 20:5149–52.

31. Moore, V.C., Strano, M.S., Haroz, E.H., et al. 2003.Individually suspended single-walled carbon nanotubes invarious surfactants. Nano Lett 3:1379–82.

32. Islam, M.F., Rojas, E., Bergey, D.M., Johnson,A.T., andYodh, A.G. 2003. High weight fraction surfactantsolubilization of single-wall carbon nanotubes in water.Nano Lett 3:269–73.

33. Strano, M.S., Moore, V.C., Miller, M.K. et al. 2003.

The role of surfactant adsorption during ultrasonication inthe dispersion of single-walled carbon nanotubes. J.Nanosci Nanotech 3:81–6.

34. Vaisman, L., Wagner, H.D., and Marom, G. 2006. The roleof surfactants in dispersion of carbon nanotubes. AdvColloid Inter Sci 128–130:37–46.

35. Cui, S., Canet, R., Derre, A., Couzi, M., and Delhaes,P. 2003. Characterization of multiwall carbon nanotubesand in�uence of surfactant in the nanocomposite processing.Carbon 41:797.

36. Shvartzman-Cohen, R., Kalisman, Y.L., Navi-Roth, E.,Yerushalmi-Rozen, R. 2004. Generic approach for dispersingsingle-walled carbon nanotubes: The strength of a weakinteraction. Langmuir 20:6085–8.

37. Hertel, T., Hagen, A., Talalaev, V. et al. 2005.Spectroscopy of single- and double-wall carbon nanotubesin different environments. Nano Lett 5:511–4.

38. Yurekli, K., Mitchell, C.A., and Krishnamootri, R.2004. Small angle neutron scattering from surfactantassisted aqueous dispersion of carbon nanotubes. J Am ChemSoc 126:9902–3.

39. Jiang, L., Gao, L., and Sun, J. 2003. Production ofaqueous colloidal dispersions of carbon nanotubes. J CollInterf Sci 260:89–94.

40. Poulin, P., Vigolo, B., and Launois, P. 2002. Films and�bers of oriented single wall nanotubes. Carbon 40:1741–9.

41. Sáfar, G.A.M., Ribeiro, H.B., and Malard, L.M. 2008.Optical study of porphyrin-doped carbon nanotubes. ChemPhys Lett 462:109–11.

42. Tan, Y. and Resasco, D.E. 2005. Dispersion ofsingle-walled carbon nanotubes of narrow diameterdistribution. J Phys Chem B 109:14454.

43. Cheng, H.K.F., Pan, Y., Sahoo, N.G. et al. 2011.Improvement in properties of multiwalled carbonnanotube/polypropylene nanocomposites through homogeneousdispersion with the aid of surfactants. J Appl Polym Sci124:1117–27.

44. Gong, X., Liu, J., Baskaran, S., Voise, R.D., andYoung, J.S. 2000. Surfactant-assisted processing of carbon

nanotube/polymer composites. Chem Mater 12:1049.

45. Sun, Z., Nicolosi, V., Rickard, D., Bergin, S.D.,Aherne, D., and Coleman, J.N. 2008. Quantitative evaluationof surfactant-stabilized single-walled carbon nanotubes:Dispersion quality and its correlation with zetapotential. J Phys Chem C 112:10692–9.

46. Fujigaya, T., Yoo, J., and Nakashima, N. 2011. A methodfor the coating of silica spheres with an ultrathin layerof pristine single-walled carbon nanotubes. Carbon49:468–76.

47. Menna, E., Negra, F.D., Prato, M. et al. 2006. Carbonnanotubes on HPLC silica microspheres. Carbon 44:1609–13.

48. Debnath, S., Cheng, Q.H., Hedderman, T.G., and ByrneH.J. 2008. An experimental study of the interaction betweensingle walled carbon nanotubes and polycyclic aromatichydrocarbons. Phys. Status Solidi B 245:1961–3.

49. Yang, K., Wang, X., Zhu, L., and Xing, B. 2006.Competitive sorption of pyrene, phenanthrene, andnaphthalene on multiwalled carbon nanotubes. Environ SciTechnol 40:5804–10.

50. Fujigaya, T. and Nakashima, N. 2008. Methodology forhomogeneous dispersion of single-walled carbon nanotubesby physical modi�cation. Polym J 40:577–89.

51. Nakashima, N. and Fujigaya, T. 2007. Fundamentals andapplications of soluble carbon nanotubes. Chem Lett36:692–7.

52. Backes, C., Hauke, F., and Hirsch, A. 2011. Thepotential of perylene bisimide derivatives for thesolubilization of carbon nanotubes and graphene. Adv Mater23:2588–601.

53. Wenseleers, W., Vlasov, I.I., Goovaerts, E.,Obraztsova, E. D., Lobach, A. S., and Bouwen, A. 2004.Ef�cient isolation and solubilization of pristinesingle-walled nanotubes in bile salt micelles. Adv FunctMater 14:1105–12.

54. Debnath, S., Cheng, Q., Hedderman, T.G., and Byrne,H.J. 2008. A study of the interaction betweensingle-walled carbon nanotubes and polycyclic aromatichydrocarbons: Toward structure−property relationships. JPhys Chem C 112:10418–22.

55. Yoo, J.T., Ozawa, H., Fujigaya, T., and Nakashima, N.2011. Evaluation of af�nity of molecules for carbonnanotubes. Nanoscale 3:2517–22.

56. Bandyopadhyaya, R., Nativ-Roth, E., Regev, O., andYerushalmi-Rozen, R. 2002. Stabilization of individualcarbon nanotubes in aqueous solutions. Nano Lett 2:25–8.

57. Pei, X., Hu, L., Liu, W., and Hao, J. 2008. Synthesisof water-soluble carbon nanotubes via surface initiatedredox polymerization and their tribological properties aswater-based lubricant additive. Eur Polym J 44:2458–64.

58. Wu, H.X., Qiu, X.Q., Cao, W.M., Lin, Y.H., Cai, R.F.,and Qian, S.X. 2007. Polymer-wrapped multiwalled carbonnanotubes synthesized via microwave-assisted in situemulsion polymerization and their optical limitingproperties. Carbon 45:2866–72.

59. Cheng, F., Imin, P., Maunders, C., Botton, G., andAdronov, A. 2008. Soluble, discrete supramolecularcomplexes of single-walled carbon nanotubes with�uorene-based conjugated polymers. Macromolecules41:2304–8.

60. Yang, L., Zhang, B., Liang, Y., Yang, B., Kong, T., andZhang, L.M. 2008. In situ synthesis of amylose/single-walled carbon nanotubes supramolecular assembly.Carbohydrate Res 343:2463–7.

61. Ago, H., Petritsch, K., Shaffer, M.S.P., Windle, A.H.,and Friend, R.H. 1999. Composites of carbon nanotubes andconjugated polymers for photovoltaic devices. Adv Mater11:1281–3.

62. Curran, S.A., Ajayan, P.M., Blau, W.J. et al. 1998. Acomposite from

63. McCarthy, B., Coleman, J.N., Czerw, R., Dalton, A.B.,Carroll, D.L., and Blau, W.J. 2001. Microscopy studies ofnanotube-conjugated polymer interactions. Synt Met121:1225–6.

64. Dalton, A.B., Coleman, C.S.J.N., McCarthy, B. et al.2000. Selective interaction of a semiconjugated organicpolymer with single-wall nanotubes. J Phys Chem B104:10012–6.

65. Cheng, F.Y., Imin, P., Maunders, C., Botton, G., and

Adronov, A. 2008. Soluble, discrete supramolecularcomplexes of single-walled carbon nanotubes with�uorene-based conjugated polymers. Macromolecules41:2304–8.

66. Imin, P., Cheng, F.Y., and Adronov, A. 2011.Supramolecular complexes of single walled carbon nanotubeswith conjugated polymers. Polym Chem 2:411–6.

67. Gurevitch, I. and Srebnik, S. J. 2008. Conformationalbehavior of polymers adsorbed on nanotubes. Chem Phys128:144901-1-8.

68. Shvartzman-Cohen, R., Nativ-Roth, E., Baskaran, E.,Levi-Kalisman, Y., Szleifer, I., and YerushalmiRozen, R.2004. Selective dispersion of single-walled carbonnanotubes in the presence of polymers: The role ofmolecular and colloidal length scales. J Am Chem Soc126:14850–7.

69. Nativ-Roth, E., Shvartzman-Cohen, R., Bounioux, C.et al. 2007. Physical adsorption of block copolymers toSWNT and MWNT: A nonwraping mechanism. Macromolecules40:3676–85.

70. Hadid, N., Kobarfard, F., Na�ssi-Varcheh, N., andAboofazeli, R. 2011. Optimization of single-walled carbonnanotube solubility by noncovalent PEGylation usingexperimental design methods. Int J Nanomed 6:737–46.

71. Niyogi, S., Hamon, M.A., and Hu, H. 2002. Chemistry ofsingle-walled carbon nanotubes. Acc Chem Res 35:1105–13.

72. Zhang, X., Sreekumar, T.V., Liu, T., and Kumar, S.2004. Properties and structure of nitric acid oxidizedsingle wall carbon nanotube �lms. J Phys Chem B108:16435–40.

73. Cho, J.W., Kim, J.W., Jung, Y.C., and Goo, N.S. 2005.Electroactive shape-memory polyurethane compositesincorporating carbon nanotubes. Macromol Rapid Commun26:412–6.

74. Georgakilas, V., Kordatos, K., Prato, M., Guldi, D.M.,Holzingger, M., and Hirsch, A. 2002. Organicfunctionalization of carbon nanotubes. J Am Chem Soc124:760–1.

75. Zhang, Y., Shi, Z., Gu, Z., and Iijima, S. 2000.Structure modi�cation of single-wall carbon nanotubes.

Carbon 38:2055–9.

76. Datsyuk, V., Kalyva, M., Papagelis, K. et al. 2008.Chemical oxidation of multiwalled carbon nanotubes. Carbon46:833–40. F r cataloging purp ses nly

77. Chen, J., Rao, A.M., Lyuksyutov, S. et al. 2001.Dissolution of full-length single-walled carbon nanotubes.J Phys Chem B 105:2525–8.

78. Thess, A., Lee, R., Nikolaev, P. et al. 1996.Crystalline ropes of metallic carbon nanotubes. Science273:483–7.

79. Mickelson, E.T., Chiang, I.W., Zimmerman, J.L. et al.1999. Solvation of �uorinated single-wall carbon nanotubesin alcohol solvents. J Phys Chem B 103:4318–22.

80. Kelly, K.F., Chiang, I.W., and Mickelson, E.T. 1999.Insight into the mechanism of sidewall functionalization ofsingle-walled nanotubes: An STM study. Chem Phys Lett313:445–50.

81. Boul, P.J., Liu, J., and Mickelson, E.T. 1999.Reversible sidewall functionalization of buckytubes. ChemPhys Lett 310:367–72.

82. Cho, E., Kim, H., Kim, C.W., and Han, S. 2006. Abinitio study on the carbon nanotube, with various degreesof functionalization. Chem Phys Lett 419:134–8.

83. Steinmetz, J., Samaille, D., Glerup, M. et al. 2004. [2+ 1] cycloaddition for cross-linking SWCNTs. Carbon42:941–7.

84. Brunetti, F.G., Herrero, M.A., Muoz, J.D.M. et al.2008. Microwave-induced multiple functionalization ofcarbon nanotubes. J Am Chem Soc 130:8094–100.

85. Chen, Y., Haddon, R.C., Fang, S. et al. 1998. Chemicalattachment of organic functional groups to singlewalledcarbon nanotube material. J Mater Res 13:2423–31.

86. Chen, J., Hamon, M.A., Hu, H. et al. 1998. Solutionproperties of single-walled carbon nanotubes. Science282:95–8.

87. Lee, W.H., Kim, S.J., Lee, W.J., Lee, J.G., Haddon,R.C., and Reucroft, P.J. 2001. X-ray photoelectronspectroscopic studies of surface modi�ed single-walled

carbon nanotube material. Appl Surf Sci 181, 121–7.

88. Kamaras, K., Itkis, M.E., Hu, H., Zhao, B., and Haddon,R.C. 2003. Covalent bond formation to a carbon nanotubemetal. Science 301, 1501.

89. Hirsch, A. 2000. Functionalization of single-walledcarbon nanotubes. Angew Chem Int Ed 41:1853.

90. Hirsch, A. and Vostrowsky, O. 2005. Functionalizationof carbon nanotubes. Top Curr Chem 245, 193–237.

91. Holzinger, M., Vostrovsky, O., Hirsch, A. et al. 2001.Sidewall functionalization of carbon nanotubes. Angew ChemInt Ed 40:4002.

92. Holzinger, M., Abraham, J., Whelan, P. et al. 2003.Functionalization of single-walled carbon nanotubes with(R-)oxycarbonyl nitrenes. J Am Chem Soc 125:8566–80.

93. Sahoo, N.G., Bao, H., Pan, Y. et al. 2011.Functionalized carbon nanomaterials as nanocarriers fordrug loading and delivery of poorly water solubleanticancer drug: A comparative study. Chem Comm 47:5235–7.

94. Riggs, J.E., Guo, Z., Carroll, D.L., and Sun, Y.P.2000. Strong luminescence of solubilized carbon nanotubes.J Am Chem Soc 122:5879–80.

95. Lin, Y., Zhou, B., Fernando, K.A.S., Liu, P., Allard,L.F., and Sun, Y.P. 2003. Polymeric carbon nanocompositesfrom carbon nanotubes functionalized with matrix polymer.Macromolecules 36:7199–204.

96. Hill, D.E,. Lin, Y., Rao, A.M., Allard, L.F., and Sun,Y.P. 2002. Functionalization of carbon nanotubes withpolystyrene. Macromolecules 35:9466–71.

97. Qin, S., Qin. D., Ford, W.T., Resasco, D.E., andHerrera, J.E. 2004. Functionalization of single-wallednanotubes with polystyrene via grafting to and graftingfrom methods. Macromolecules 37:752–7.

98. Jung, Y.C., Sahoo, N.G., and Cho, J.W. 2006. Polymericnanocomposites of polyurethane block copolymers andfunctionalized multi-walled carbon nanotubes ascrosslinkers. Macromol Rapid Commun 27:126–31.

99. Kuan, H.C., Ma, C.C.M., Chang, W.P., Yuen, S.M., Wu,H.H., and Lee, T.M. 2005. Synthesis, thermal, mechanical

and rheological properties of multiwall carbonnanotube/waterborne polyurethane nanocomposite. Compos SciTechnol 65:1703–10.

100. Gohier, A.E., Nekelson, F., Helezen, M. et al. 2011.Tunable grafting of functional polymers onto carbonnanotubes using diazonium chemistry in aqueous media. JMater Chem 21:4615.

101. Qin, S., Qin, D., Ford, W.T. et al. 2004.Solubilization and puri�cation of single-wall carbonnanotubes in water by in situ radical polymerization ofsodium 4-styrenesulfonate. Macromolecules 37:3965–7.

102. Koshio, A., Yudasaka, M., Zhang, M., and Iijima, S.2001. A simple way to chemically react single-wall carbonnanotubes with organic materials using ultrasonication.Nano Lett 1:361–3.

103. Qu, L.W., Lin, Y., Hill, D.E. et al. 2004.Polyimidefunctionalized carbon nanotubes: Synthesis anddispersion in nanocomposite �lms. Macromolecules37:6055–60.

104. Park, C., Ounaies, Z., Watson, K.A. et al. 2002.Dispersion of single wall carbon nanotubes by in situpolymerization under sonication. Chem Phys Lett 364:303–8.

105. Lou, X.D., Detrembleur, C., Pagnoulle, C. et al. 2004Surface modi�cation of multiwalled carbon nanotubes bypoly(2-vinylpyridine): Dispersion, selective deposition,and decoration of the nanotubes. Adv Mater 16:2123–7.

106. Riggs, J.E., Walker, D.B., Carroll, D.L., Sun. Y.-P.2000. Optical limiting properties of suspended andsolubilized carbon nanotubes. J Phys Chem B 104:7071–6.

107. Sun, Y.-P., Huang, W., Lin, Y. et al. 2001. Solubledendron-functionalized carbon nanotubes: Preparation,characterization, and properties. Chem Mater 13:2864–69.

108. Sano, M., Kamino, A., and Shinkai, S. 2001.Construction of carbon nanotube stars with dendrimers.Angew Chem Int Ed 40:4661–3.

109. Cao, L., Yang, W., Yang, J., Wang, C., and Fu. S.2004. Hyperbranched poly(amidoamine)-modi�ed multi-walledcarbon nanotubes via grafting-from method. Chem Lett33:490–1.

110. Chen, S.M., Wu, G.Z., Liu, Y.D., and Long, D.W. 2006.Preparation of poly(acrylic acid) grafted multiwalledcarbon nanotubes by a two-step irradiation technique.Macromolecules 39:330–34.

111. Liu, Y.Q., Yao, Z.L., and Adronov, A. 2005.Functionalization of single-walled carbon nanotubes withwell-de�ned polymers by radical coupling. Macromolecules38:1172–79.

112. Guo, G.Q., Yang, D., Wang, C.C., and Yang, S. 2006.Fishing polymer brushes on single-walled carbon nanotubesby in-situ free radical polymerization in a poor solvent.Macromolecules 39:9035–40.

113. Peng, M., Liao, Z., Zhu, Z., and Guo, H. 2010. Asimple polymerizable polysoap greatly enhances thegrafting ef�ciency of the “grafting-to” functionalizationof multiwalled carbon nanotubes. Macromolecules43:9635–44.

114. Sakellariou, G., Ji, H., Mays, J. W., and Baskaran, D.2008. Enhanced polymer grafting from multiwalled carbonnanotubes through living anionic surface-initiatedpolymerization. Chem Mater 20:6217–30.

115. Priftis, D., Petzetakis, N., and Sakellariou, G. 2009.Surface-initiated titanium-mediated coordinationpolymerization from catalyst-functionalized single andmultiwalled carbon nanotubes. Macromolecules 42:3340–6.

116. Lee, R.S., Chen, W.S., and Lin, J.H. 2011.Polymer-grafted multi-walled carbon nanotubes throughsurface-initiated ring-opening polymerization and clickreaction. Polymer 52:2180–8.

117. Yao, Y., Li, W., Wang, S., Yan, D., and Chen, X. 2006.Polypeptide modi�cation of multiwalled carbon nanotubes bya graft-from approach. Macromol Rapid Commun 27:2019.

118. Le, A.P., Huang, T.M., Chen, P.T., and Yang, A.C.M.2001. Synthesis and optoelectronic behavior of conjugatedpolymer poly(3-hexylthiophene) grafted on multiwalledcarbon nanotubes. J Appl Sci Part B Polym Phys 49:581–90.

119. Chen, X.H., Chen, X., Lin, M., Zhong, W., Chen, X.,and Chen, Z. 2007. Functionalized multi-walled carbonnanotubes prepared by in situ polycondensation ofpolyurethane. Macrom Chem Phys 208:964–72.

120. Qin, S., Qin, D., Ford, W.T., Resasco, D.E., andHerrera, J.E. 2004. Polymer brushes on single-walled carbonnanotubes by atom transfer radical polymerization ofn-butyl methacrylate. J Am Chem Soc 126:170–6.

121. Kong, H., Gao, C., and Yan, D. 2004. Controlledfunctionalization of multiwalled carbon nanotubes by insitu atom transfer radical polymerization. J Am Chem Soc126:412–3.

122. Qu, L., Veca, L.M., Lin, Y. et al. 2005. Solublenylon-functionalized carbon nanotubes from anionicringopening polymerization from nanotube surface.Macromolecules 38:10328–31.

123. Yang, M., Gao, Y., Li, H., and Adronov, A. 2007.Functionalization of multiwalled carbon nanotubes withpolyamide 6 by anionic ring-opening polymerization. Carbon45:2327–33.

124. Park, S.J., Cho, M.S., Lim, S.T., Cho, H.J., and Jhon,M.S. 2003. Synthesis and dispersion characteristics ofmulti-walled carbon nanotube composites with poly(methylmethacrylate) prepared by in-situ bulk polymerization.Macromol Rapid Commun 24:1070–3.

125. Yao, Z., Braidy, N., Botton, G.A., and Adronov, A.2003. Polymerization from the surface of single-walledcarbon nanotubes—preparation and characterization ofnanocomposites. J Am Chem Soc 125:16015–24.

126. Baskaran, D., Mays, J.W., and Bratcher, M.S. 2004.polymer-grafted multiwalled carbon nanotubes throughsurface-initiated polymerization. Angew Chem Int Ed43:2138–42.

127. Kong, H., Gao, C., and Yan, D. 2004. Functionalizationof multiwalled carbon nanotubes by atom transfer radicalpolymerization and defunctionalization of the products.Macromolecules 37:4022–30.

128. Kong, H., Luo, P., Gao, C., and Yan, D. 2005.Polyelectrolyte-functionalized multiwalled carbonnanotubes: Preparation, characterization and layer-by-layerself-assembly. Polymer 46:2472–85.

129. Kong, H., Gao, C., and Yan, D. 2004. Constructingamphiphilic polymer brushes on the convex surfaces ofmulti-walled carbon nanotubes by in situ atom transferradical polymerization. J Mater Chem 14:1401–5. For

catalogi g purposes only

130. Kong, H., Li, W., Gao C. et al. 2004.Poly(N-isopropylacrylamide)-coated carbon nanotubes:Temperaturesensitive molecular nanohybrids in water.Macromolecules 37:6683–6.

131. Qin, S., Qin, D., Ford, W.T., Herrera, J.E., andResasco, D.E. 2004. Grafting of poly(4-vinylpyridine) tosingle-walled carbon nanotubes and assembly of multilayer�lms. Macromolecules 37:9963–7.

132. Wu, W., Zhang, S., Li, Y. et al. 2003. PVK-modi�edsinglewalled carbon nanotubes with effective photoinducedelectron transfer. Macromolecules 36:6286–8.

133. Li, C., Pang, X.J., and Qu, M.Z. 2005. Fabrication andcharacterization of polycarbonate/carbon nanotubescomposites. Compos Part A 37:1485–9.

134. Sahoo, N.G., Jung, Y.C., Yoo. H.J., and Cho, J.W.2006. Effect of functionalized carbon nanotubes onmolecular interaction and properties of polyurethanecomposites. Macromol Chem Phys 207:1773–80.

135. Chen, W. and Tao, X. 2005. Self-organizing alignmentofcarbon nanotubes in thermoplastic polyurethane. MacromoRapid Commun 26:1763.

136. Vigolo, B., Penicaud, A., Coulon, C. et al. 2000.Macroscopic �bers and ribbons of oriented carbon nanotubes.Science 290:1331–4.

137. Saran, N., Parikh, K., Suh, D.S., Munoz, E., Kolla,H., and Manhor, S.K. 2004. Fabrication and characterizationof thin �lms of single-walled carbon nanotubes bundles on�exible plastic substrates. J Am Chem Soc 126:4462–3.

138. Moniruzzaman, M. and Winey, K.I. 2006. Polymernanocomposites containing carbon nanotubes. 39:5194–205.

139. Bryning, M.B., Milkie, D.E., Islam, M.F., Kikkawa,J.M., and Yodh, A.G. 2005. Thermal conductivity andinterfacial resistance in single-wall carbon nanotube epoxycomposites. Appl Phys Lett 87:161909/1–3.

140. De la Chapelle, M.L., Stephan, C., Nguyen, T.P. et al. 1999. Raman characterization of singlewalled carbonnanotubes and PMMA-nanotubes composites. Synth Met103:2510–2.

141. Benoit, J.M., Corrage, B., Lefrant, S., Blau, W.J.,Bernier, P., and Chauvet, O. 2001. Transport properties ofPMMA-carbon nanotubes composites. Synth Met 121:1215–6.

142. Du, F., Fischer, J.E., and Winey, K.I. 2003.Coagulation method for preparing singlewalled carbonnanotube/poly(methyl methacrylate) composites and theirmodulus, electrical conductivity, and thermal stability. JPolym Sci B 41:3333–8.

143. Qian, D., Dickey, E. C., Andrews, R., and Rantell, T.2000. Load transfer and deformation mechanisms in carbonnanotube-polystyrene composites. Appl Phys Lett 76:2868–70.

144. Kota, A.K., Cipriano, B.H., Duesterberg, M.K. et al.2007. Electrical and reheological percolation inpolystyre/MWCNT nanocomposites. Macromolecules 40:7400–6.

145. Song, Y.S. and Youn, J.R. 2005. In�uence of dispersionstates of carbon nanotubes on physical properties of epoxynanocomposites. Carbon 43:1378–85.

146. Shaffer, M.S.P. and Windle, A.H. 1999. Fabrication andcharacterization of carbon nanotube/poly(vinyl alcohol)composites. Adv Mater 11:937–41.

147. Yang, J., Hu, J., Wang, C., Qin, Y., and Guo, Z. 2004.Fabrication and characterization of soluble multiwalledcarbon nanotubes reinforced P(MMA-co-EMA) composites.Macromol Mat Eng 289:828–32.

148. Chae, H.G., Sreekumar, T.V., Uchida, T., and Kumar, S.2005. A comparison of reinforcement ef�ciency of varioustypes of carbon nanotubes in polyacrylonitrile �ber.Polymer 46:10925–35.

149. Bin, Y., Kitanaka, M., Zhu, D., and Matsuo, M. 2003.Development of highly oriented polyethylene �lled withaligned carbon nanotubes by gelation/crystallization fromsolutions. Macromolecules 36:6213–9.

150. Andrews, R., Jacques, D., Minot, M., and Rantell, T.2002. Fabrication of carbon multiwall nanotube/ polymercomposites by shear mixing. Micromol Mater Eng 287:395–403.

151. Sahoo, N.G., Cheng, H.K.F., Bao, H., Li, L., Chan,S.H., and Zhao, J. 2011. Nitrophenyl functionalization ofcarbon nanotubes and its effect on properties of MWCNT/LCPnanocomposites. Macromol Res 19:660–7.

152. McNally, T., Potschkeb, P., Halley et al. 2005.Polyethylene multiwalled carbon nanotube composites.Polymer 46:8222–32.

153. Sahoo, N.G., Cheng, H.K.F., Tan, Q.H., Li, L., Chan,S.H., and Zhao, J. 2009. Effect of carbon nanotubes andprocessing methods on the properties of carbonnanotube/polypropylene composites. J Nanosci Nanotech9:5910–9.

154. Bellayer, S., Gilman, J.W., Eidelman, N. et al. 2005.Preparation of homogeneously dispersed multiwalled carbonnanotube/polystyrene nanocomposites via melt extrusionusing trialkyl imidazolium compatibilizer. Adv Funct Mater15:910–6.

155. Yang, B.X., Pramoda, K.P., Xu, G.Q., and Goh, S.H.2007. Mechanical reinforcement of polyethylene usingpolyethylene-grafted multiwalled carbon nanotubes. AdvFunct Mater 17:2062–9.

156. Prashantha, K., Soulestin, J., Lacrampe, M.F.,Krawczak, P., Dupin, G., and Claes, M. 2009.Masterbatchbased multi-walled carbon nanotube �lledpolypropylene nanocomposites: Assessment of rheologicaland mechanical properties. Compos Sci Technol 69:1756–63.

157. Pan, Y., Li, L., Chan, S.W., and Zhao, J. 2010.Correlation between dispersion state and electricalconductivity of MWCNTs/PP composites prepared by meltblending. Compos Part A 41:419–26.

158. Bocchini, S., Frache, A., Camino, G., and Claes, M.2007. Polyethylene thermal oxidative stabilisation incarbon nanotubes based nanocomposites. Eur Polym J43:3222–35.

159. Tang, W., Santare, M.H., and Advani, S.G. 2003. Meltprocessing and mechanical property characterization ofmulti-walled carbon nanotube/high density polyethylene(MWNT/HDPE) composite �lms. Carbon 41:2779–85.

160. Pötschke, P., Bhattacharyya, A.R., and Janke, A. 2004.Melt mixing of polycarbonate with multiwalled carbonnanotubes: Microscopic studies on the state of dispersion.Eur Polym J 40:137–48.

161. Zeng, J., Saltysiak, B., Johnson, W.S., Schiraldi,D.A., and Kumar, S. 2004. Processing and properties of

poly(methyl methacrylate)/carbon nano �ber composites.Compos Part B 35:173–8.

162. Gorga, R.E. and Cohen, R.E. 2004. Toughnessenhancements in poly(methyl methacrylate) by addition oforiented multiwall carbon nanotubes. J Polym Sci Part BPolym Phys 42:2690–702.

163. Siochi, E.J., Working, D.C., Park, C. et al. 2004.Melt processing of SWNT-polyimide nanocomposite �bres.Compos Part B 35:439–46.

164. Meincke, O., Kaempfer, D., Weickmann, H., Friedrich,C., Vathauer, M., and Warth, H. 2004. Mechanicalproperties and electrical conductivity of carbon-nanotube�lled polyamide-6 and its blends withacrylonitrile/butadiene/styrene. Polymer 45:739–48.

165. Ajayan, P.M., Schadler, L. S., Giannaris, C., andRubio, A. 2000. Single-walled carbon nanotube-polymercomposites: Strength and weakness. Adv Mater 12:750.

166. Zhu, J., Peng, H., Rodriguez-Macias, F. et al. 2004.Reinforcing epoxy polymer composites through covalentintegration of functionalized nanotubes. Adv Funct Mater14:643–8.

167. Ma, P.C., Mo, S.Y., Tang, B.Z., Kim, J.K. 2010.Dispersion, interfacial interaction and re-agglomerationof functionalized carbon nanotubes in epoxy composites.Carbon 48:1824–34.

168. Putz, K.W., Mitchell, C.A., Krishnamoorti, R., andGreen, P.F. 2004. Elastic modulus of single-walled carbonnanotube/poly(methyl methacrylate) nanocomposites. J PolymSci Part B Polym Phys 42:2286–93.

169. Sung, J.H., Kim, H.S., Jin, H.J., Choi, H.J., andChin, I.J. 2004. Nano�brous membranes prepared bymultiwalled carbon nanotube/poly(methyl methacrylate)composites. Macromolecules 37:9899–902.

170. Jia, Z., Wang, Z., Xu, C. et al. 1999. Study onpoly(methyl methacrylate): Carbon nanotube composites.Mater Sci Eng A 271:395–400.

171. Xia, H., Song, M., Jin, J., and Chen, L. 2006.Poly(propylene glycol)-grafted multiwalled carbon nanotubepolyurethane. Macromol Chem Phys 27:1945–52.

172. Karim, M.R., Lee, C.J., Park, Y.T., and Lee, M.S.2005. SWNTs coated by conducting polyaniline: Synthesisand modi�ed properties. Synth Met 151:131–5.

173. Long, Y., Chen, Z., Zhang, X., Zhang, J., and Liu, Z.2004. Electrical properties of multiwalled carbonnanotube/polypyrrole nanocables: Percolationdominatedconductivity. J Phys D Appl Phys 37:1965–69.

174. Sahoo, N.G., Jung, Y.C., So, H.H., and Cho, J.W. 2007.Polypyrrole coated carbon nanotubes: Synthesis,characterization, and enhanced electrical properties. SynthMet 157:179–37.

175. Mamedov, A.A., Kotov, N.A., Prato, M., Guldi, D.M.,Wicksted, J.P., and Hirsch., A. 2002. Molecular design ofstrong single-wall carbon nanotube/polyelectrolytemultilayer composites. Nat Mater 190–94.

176. Xue, W., Liu, Y., and Cui. T. 2006. High-mobilitytransistors based on nanoassembled carbon nanotubesemiconducting layer and SiO 2 nanoparticle dielectriclayer. Appl Phys Lett 89:163512.

177. Feng, Q.P., Yang, J.P., Fu, S.Y., and Mai, Y.W. 2010.Synthesis of carbon nanotube/epoxy composite �lms with ahigh nanotube loading by a mixed-curing-agent assistedlayer-by-layer method and their electrical conductivity.Carbon 48:2057–62.

178. Shim, B.S., Tang, Z.Y., Morabito, M.P., Agarwal, A.,Hong, H.P., and Kotov, N.A. 2007. Integration ofconductivity transparency, and mechanical strength intohighly homogeneous layer-by-layer composites ofsingle-walled carbon nanotubes for optoelectronics. ChemMater 19:5467.

179. Grossiord, N., Kivit, P.J.J., Loos, J. et al. 2008. Onthe in�uence of the processing conditions on theperformance of electrically conductive carbonnanotube/polymer nanocomposites. Polymer 49:2866–72.

180. Regev, O., Elkati, P.N.B., Loos, J., and Koning, C.E.2004. Preparation of conductive nanotube–polymercomposites using latex technology. Adv Mater 16:248–51.

181. Grunlan, J.C., Mehrabi, A.R., Bannon, M.V., and Bahr,J.L. 2004. Water-based single-walled-nanotube�lled polymercomposite with an exceptionally low percolation threshold.Adv Mater 16:150–3. For c talogin purp ses only

182. Grossiord, N., Loos, J., and Koning, C.E. 2005.Strategies for dispersing carbon nanotubes in highlyviscous polymers. J Mater Chem 15:2349–52.

183. Yu, J., Grossiord, N., Koning, C.E., and Loos, J.2007. Characterization of conductive multiwall carbonnanotube/polystyrene composites prepared by latextechnology. Carbon 45:618–23.

184. Hermant, M.C., Klupperman, B., Kyrylyuk, A.V., Schoot,P.V.D., and Koning, C.E. 2009. Lowering the percolationthreshold of single-walled carbon nanotubes usingpolystyrene/poly(3,4-ethylenedioxythiophene): Poly(styrenesulfonate) blends. Soft Mater 5:878–85.

185. O’Connor, I., De, S., Coleman, J.N., Gun’ko, and Y.K.2009. Development of transparent, conducting composites bysurface in�ltration of nanotubes into commercial polymer�lms. Carbon 47:1983–8.

186. Chung, T.S. 2001. Thermotropic Liquid CrystalPolymers: Thin Film Polymerization, Characterization,Blends, and Applications. London: CRC Press.

187. Demus, D., Gray, G.W., Speiss, H.W., Goodby, J.W., andVill, V. 1998. Handbook of Liquid Crystals. Weinheim:Wiley-VCH.

188. Park, S.K., Kim, S.H., and Hwang, J.T. 2008.Carboxylated multiwall carbon nanotube-reinforcedthermotropic liquid crystalline polymer nanocomposites. JAppl Polym Sci 109:388.

189. Kim, J.Y., Kim, D.K., and Kim, S.H. 2009. Effect ofmodi�ed carbon nanotube on physical properties ofthermotropic liquid crystal polyester nanocomposites. EurPolym J 45:316–24.

190. Choi, H.J., Kim, J.Y., and Kim, S.H. 2009. Multi-wallcarbon nanotube reinforced thermotropic liquid crystalcopolyester nanocomposites. Mol Cryst Liq Cryst510:300/1434–311/1445.

191. Sahoo, N.G., Cheng, H.K.F., Bao, H., Pan, Y., Li, L.,and Chan, S.H. 2011. Covalent functionalization of carbonnanotubes for ultimate interfacial adhesion to liquidcrystalline polymer. Soft Matter 7:9505–14.

192. Roy, S., Sahoo, N.G., Cheng, H.K.F., Das, C.K., Li,

L., and Chan, S.H. 2011. Molecular interaction andproperties of poly (ether ether ketone)/liquid crystallinepolymer blends incorporated with functionalized farbonnanotubes. J Nanosci Nanotech 11:10408–16.

193. Roy, S., Sahoo, N.G., Das, C.K., Li, L., and Chan,S.H. 2009. Improvement of properties ofpolyetherimide/liquid crystalline polymer blends in thepresence of functionalized carbon nanotubes. J NanosciNanotech 9:1928–34.

194. Nayak, G.C., Rajasekar, R., and Das, C.K. 2010. Effectof SiC coated MWCNTs on the thermal and mechanicalproperties of PEI/LCP blend. Composites: Part A 41:1662–7.

195. Nayak, G.C., Rajasekar, R., and Das, C.K. 2011. Effectof modi�ed MWCNT on the properties of PPO/ LCP blend. JMater Sci 46:2050–7.

196. Zhu, J., Kim, J., Peng, H., Margrave, J.L.,Khabashesku, V.N., and Barrera, E.V. 2003. Improving thedispersion and integration of single-walled carbonnanotubes in epoxy composites through functionalization.Nano Lett 3:1107–13.

197. de Villoria, R.G., Miravete, A., Cuartero, J.,Chiminelli, A., and Tolosana, N. 2006. Mechanicalproperties of SWNT/epoxy composites using two differentcuring cycles. Composites Part B 37:273–7.

198. Valentini, L., Puglia, D., Carniato, F., Boccaleri,E., Marchese, L., and Kenny, J. M. 2008. Use of plasma�uorinated single-walled carbon nanotubes for thepreparation of nanocomposites with epoxy matrix. Comp SciTechnol 68:1008–14.

199. Sun, L., Warren, G.L., O’Reilly, J.Y., Everett, W.N.,Lee, S.M., Davis, D., Lagoudas, D., and Sue, H.J. 2008.Mechanical properties of surface-functionalized SWCNT/epoxycomposites. Carbon 46:320–8.

200. Che, J., Yuan, W., Jiang, G., Dai, J., Lim, S. Y., andChan-Park, M. B. 2009. Epoxy composite �bers reinforcedwith aligned single-walled carbon nanotubes functionalizedwith generation 0−2 dendritic poly(amidoamine). Chem Mater21:1471–9.

201. Chen, Z.-K., Yang, J.-P., Ni, Q.-Q., Fu, S.-Y., andHuang, Y.-G. 2009. Reinforcement of epoxy resins withmulti-walled carbon nanotubes for enhancing cryogenic

mechanical properties. Polymer 50:4753–9.

202. Tseng, C.-H., Wang, C.-C., and Chen, C.-Y. 2006.Functionalizing carbon nanotubes by plasma modi�cation forthe preparation of covalent-integrated epoxy composites.Chem. Mater. 19:308–15.

203. Breton, Y., Désarmot, G., Salvetat, J.P., Delpeux, S.,Sinturel, C., Béguin, F., and Bonnamy, S. 2004. Mechanicalproperties of multiwall carbon nanotubes/epoxy composites:In�uence of network morphology. Carbon 42:1027–30.

204. Guo, P., Chen, X., Gao, X., Song, H., and Shen, H.2007. Fabrication and mechanical properties ofwelldispersed multiwalled carbon nanotubes/epoxycomposites. Comp Sci Technol 67:3331–7.

205. Omidi, M., Rokni D.T.H., Milani, A.S., Seethaler,R.J., and Arasteh, R. 2010. Prediction of the mechanicalcharacteristics of multi-walled carbon nanotube/epoxycomposites using a new form of the rule of mixtures.Carbon 48:3218–28.

206. Gojny, F.H., Wichmann, M.H.G., Köpke, U., Fiedler, B.,and Schulte, K. 2004. Carbon nanotube-reinforcedepoxy-composites: Enhanced stiffness and fracture toughnessat low nanotube content. Comp Sci Technol 64:2363–71.

207. Gojny, F.H., Wichmann, M.G., Fiedler, B., and Schulte,K. 2005. In�uence of different carbon nanotubes on themechanical properties of epoxy matrix composites—Acomparative study. Comp Sci Technol 65:2300–13.

208. Allaoui, A., Bai, S., Cheng, H.M., and Bai, J.B. 2002.Mechanical and electrical properties of a MWNT/ epoxycomposite. Comp Sci Technol 62:1993–8.

209. Bai, J. 2003. Evidence of the reinforcement role ofchemical vapour deposition multi-walled carbon nanotubesin a polymer matrix. Carbon 41:1325–8.

210. Ma, P. C., Kim, J.-K., and Tang, B.Z. 2007. Effects ofsilane functionalization on the properties of carbonnanotube/epoxy nanocomposites. Comp Sci Technol 67:2965–72.

211. Hernández-Pérez, A., Avilés, F., May-Pat, A.,Valadez-González, A., Herrera-Franco, P.J., andBartoloPérez, P. 2008. Effective properties of multiwalledcarbon nanotube/epoxy composites using two different tubes.Comp Sci Technol 68:1422–31.

212. Martone, A., Faiella, G., Antonucci, V., Giordano, M.,and Zarrelli, M. 2011. The effect of the aspect ratio ofcarbon nanotubes on their effective reinforcement modulusin an epoxy matrix. Comp Sci Technol 71:1117–23.

213. Yuan, W., Feng, J., Judeh, Z., Dai, J., and Chan-Park,M. B. 2010. Use of polyimide-graft-bisphenol a diglycerylacrylate as a reactive noncovalent dispersant ofsingle-walled carbon nanotubes for reinforcement of cyanateester/epoxy composite. Chem Mater 22:6542–54.

214. Wang, Q., Dai, J., Li, W., Wei, Z., and Jiang, J.2008. The effects of CNT alignment on electricalconductivity and mechanical properties of SWNT/epoxynanocomposites. Comp Sci Technol 68:1644–8.

215. Ma, W., Liu, L., Zhang, Z. et al. 2009. High-strengthcomposite �bers: Realizing true potential of carbonnanotubes in polymer matrix through continuous reticulatearchitecture and molecular level couplings. Nano Lett9:2855–61.

216. Mora, R.J., Vilatela, J.J., and Windle, A.H. 2009.Properties of composites of carbon nanotube �bres. Comp.Sci. Technol. 69:1558–63.

217. Baughman, R.H., Zakhidov, A.A., and de Heer, W.A.2002. Carbon nanotubes—the route toward applications.Science 297:787–92.

218. Ericson, L.M., Fan, H., Peng, H.Q. et al. 2004.Macroscopic, neat, single-walled carbon nanotubes �bers.Science 305:1447–50.

219. Sreekumar, T.V., Liu, T., Kumar, S., Ericson, L.M.,Hauge, R.H., and Smalley, R.E. 2003. Single-walled carbonnanotube �lms. Chem Mater 15:175–8.

220. Sandler, J., Shaffer, M.S.P., Prasse, T., Bauhofer,W., Schulte, K., and Windle, A.H. 1999. Development of adispersion process for carbon nanotubes in an epoxy matrixand the resulting electric properties. Polymer 40:5967–71.

221. Kymakis, E., Alexandou, I., and Amaratunga, G.A.J.2002. Single-walled carbon nanotube-polymer composites:Electrical, optical and structural investigation. Syn Met127:59–62.

222. Benoit, J.M., Benoit, C., Lefrant, S., Bernier, P.,

and Chauvet, O. 2002. Electric transport properties andpercolation in carbon nanotube/PMMA composites. Mater EsSoc Symp Proc 706:Z3.28.1-Z3.28.6.

223. Sandler, J.K.W., Kirk, J.E., Kinloch, I.A., Shaffer,M.S.P., and Windle, A.H. 2003. Ultra-low electricalpercolation threshold in carbon-nanotube-epoxy composites.Polymer 44:5893–9.

224. Li, H.C., Lu, S.Y., Syue, S.H., Hsu, W.K., and Chang,S.C. 2008. Conductivity enhancement of carbon nanotubecomposites by electrolyte addition. Appl Phys Lett93:033104/1–3.

225. Kim, B., Lee, J., and Yu. I.S. 2003. Electricalproperties of single-wall carbon nanotube and epoxycomposites. J Appl Phys 94:6724–8.

226. Ounaies, Z., Park, C., Wise, K.E., and Siochi, E.J.,Harrison J.S. 2003. Electrical properties of single wallcarbon nanotube reinforced polyimide composites. Comp SciTech 63:1637–46.

227. Li, J., Ma, P.C., Chow, W.S., To, C.K., Tang, B.Z.,and Kim, J.K. 2007. Correlations between percolationthreshold, dispersion state, and aspect ratio of carbonnanotubes. Adv Funct Mater 17:3207–15.

228. Choi, E.S., Brooks, J.S., Eaton, D.L. et al. 2003.Enhancement of thermal and electrical properties of carbonnanotube polymer composites by magnetic �eld processing. JAppl Phys 94:6034–9.

229. Fangming, D., Fisher, J.E., and Winey, K.I. 2003.Coagulation method for preparing single-walled carbonnanontube/poly(methyl methacrylate) composites and theirmodulus, electrical conductivity and thermal stability. JPolym Sci Part B Polym Phys 41:3333–8.

230. Bai, J.B. and Allaoui, A. 2003. Effect of the lengthand the aggregate size of MWNTs on the improvementef�ciency of the mechanical and electrical properties ofnanocomposites-experimental investigation. Compos Part A34:689–94. For cataloging pu po es only

231. Bryning, M.B., Islam, M.F., Kikkawa, J.M., and Yodh,A.G. 2005. Very low conductivity threshold in bulkisotropic single-walled carbon nanotube-epoxy composites.Adv Mater 17:1186–91.

232. Pan, Y., Cheng, H.K.F., Li, L., Chan, S.H., Zhao, J.,and Juay, Y.K. 2010. Annealing induced electricalconductivity jump of multi-walled carbonnanotube/polypropylene composites and in�uence of molecularweight of polypropylene. J Polym Sci Part B Polym Phys48:2238–47.

233. Ramasubramaniam, R., Chen, J., and Liu, H. 2003.Homogeneous carbon nanotube/polymer composites forelectrical applications. Appl Phys Lett 83:2929–30.

234. Tamburri, E., Orlanducci, S., Terranova, M.L. et al.2005. Modulation of electrical properties in singlewalledcarbon nanotube/conducting polymer composites. Carbon43:1213–21.

235. Li, S., Qin, Y., Shi, J., Guo, Z.X., Li, Y., and Zhu,D. 2005. Electrical properties of soluble carbonnanotube/polymer composite �lms. Chem Mater 17:130–5.

236. Cheng, H.K.F., Sahoo, N.G., Li, L., Chan, S.H., andZhao, J. 2010. Complementary effects of multiwalled carbonnanotubes and conductive carbon black on polyamide 6. JPolym Sci Part B Polym Phys 48:1203–12.

237. Yang, Y.L., Gupta, M.C., Dudley, K.L., and Lawrence,R.W. 2005. A comparative study of EMI shielding propertiesof carbon nano�ber and multi-walled carbon nanotube �lledpolymer composites. J Nanosci Nanotechnol 5:927–31.

238. Kim, S.H., Jang, S.H., Byun, S.W. et al. 2003.Electrical properties and EMI shielding characteristics ofpolypyrrole–nylon 6 composite fabrics. J Appl Polym Sci87:1969–74.

239. Li, N., Huang, Y., Du, F. et al. 2006. Electromagneticinterference (EMI) shielding of single-walled carbonnanotube epoxy composites. Nano Lett 6:1141–5.

240. Sohi, N.J.S., Rahaman, M., and Khastgir, D. 2011.Dielectric Property and electromagnetic interferenceshielding effectiveness of ethylene vinyl acetate-basedconductive composites: Effect of different type of carbon�llers. Polym Compos 32(7):1148–54.

241. Joo, J. and Epstein, A.J. 1994.Electromagnetic-radiation shielding by intrinsicallyconducting polymers. Appl Phys Lett 65:2278–80.

242. Joo, J. and Lee, C.Y. 2000. High frequency

electromagnetic interference shielding response of mixturesand multilayer �lms based on conducting polymers. J ApplPhys 88:513–8.

243. Bryning, M.B., Islam, M.F., Kikkawa, J.M., and Yodh,A.G. 2005. Very low conductivity threshold in bulkisotropic single-walled carbon nanotube–epoxy composites.Adv Mater 17:1186–91.

244. Chung, D.D.L. 2001. Electromagnetic interferenceshielding effectiveness of carbon materials. Carbon39:279–85.

245. Huang, Y., Li, N., Ma, Y. et al. 2007. The in�uence ofsingle-walled carbon nanotube structure on theelectromagnetic interference shielding ef�ciency of itsepoxy composites. Carbon 45:1614–21.

246. Pande, S., Singh, B.P., Mathur, R.B., Dhami, T.L.,Saini, P., and Dhawan, S.K. 2010. Improved electromagneticinterference shielding properties of MWCNT–PMMA compositesusing layered structures. Nanoscale Res Lett 4:327–34.

247. Yang, Y., Gupta, M.C., and Dudley, K.L. 2007. Towardscost-ef�cient EMI shielding materials using carbonnanostructure-based composites. Nanotechnology 18:345701–4.

248. Liu, Z., Bai, G., Haung, Y. et al. 2007. Re�ection andabsorption contribution to the electromagneticinterference shielding of single-walled carbonnanotube/polyurethane composites. Carbon 45:821–7.

249. Gupta, A. and Choudhary, V. 2011. Electromagneticinterference shielding behavior of poly(trimethyleneterephthalate)/multi-walled carbon nanotube composites.Compos Sci Technol 71:1563–8.

250. Saini, P., Choudhary, V., Singh, B.P., Mathur, R.B.,and Dhawan, S.K. 2009. Polyaniline–MWCNT nanocomposites formicrowave absorption and EMI shielding. Macromol Chem Phys113:919–26.

251. Al-Saleh, M.H. and Sundararaj, U. 2009.Electromagnetic shielding mechanism of CNT/polymercomposites. Carbon 47:1738–46.

252. O’Donnell, S.E., Sprong, K.R., and Haltll, B.M. 2004.Potential impact of carbon nanotube reinforced polymercomposite on commercial aircraft performance and economics.In: AIAA 4th Aviation Technology. Integration and

Operations (ATIO) Forum. Sep. 22–22, 2004; Chicago,Illinois.

253. Ren, Y., Li, F., Cheng, H.M., and Liao, K. 2003.Fatigue behaviour of unidirectional single-walled carbonnanotube reinforced epoxy composite under tensile load. AdvCompos Lett 12:19–20.

254. Gu, A., Liang, G., Liang, D., and Ni, M. 2007.Bismaleimide/carbon nanotube hybrids for potentialaerospace application: I. Static and dynamic mechanicalproperties. Polym Adv Technol 18:835–40.

255. Kireitseu, M., Hui, D., and Tomlinson, G. 2008.Advanced shock-resistant and vibration damping ofnanoparticle-reinforced composite material. Compos Part B39:128–38.

256. Joshi, S.C. and Dikshit, V. 2011. Enhancinginterlaminar fracture characteristics of woven CFRP prepregcomposites through CNT dispersion. J Compos Mater, DOI:10.1177/0021998311410472.

257. Wu, H.P., Wu, X.J., Ge, M.Y., Zhang, G.Q., Wang, Y.W.,and Jiang, J. 2007. Properties investigation onisotropical conductive adhesives �lled with silver coatedcarbon nanotubes. Compos Sci Technol 67:1182–6.

258. Li, J. and Lump, J.K. 2007. Carbon nanotube �lledconductive adhesives for aerospace applications. IEEEAerospace Conference 1–6.

259. Rosca, I.D. and Hoa, S.V. 2011, Method for reducingcontact resistivity of carbon nanotube-containing epoxyadhesives for aerospace applications. Compos Sci Technol71:95–100.

260. Bierwagen, G.P., and Tallman, D.E. 2001. Choice andmeasurement of crucial aircraft coatings systemproperties. Prog Org Coat 41:201–16.

261. Smith, J.G., Connell, J.W., Delozier, D.M. et al.2004. Space durable polymer/carbon nanotube �lms forelectrostatic charge mitigation. Polymer 45:825–36.

262. Zhao, W., Li, M., and Peng, H.X. 2010. FunctionalizedMWNT-doped thermoplastic polyurethane nanocomposites foraerospace coating applications. Macromol Mater Eng295:838–45.

263. Ding, K., Hu, B., Xie, Y. et al. 2009. A simple routeto coat mesoporous SiO 2 layer on carbon nanotubes. JMater Chem 19:3725–31.

264. Das, D. and Das, P.K. 2009. Superior activity ofstructurally deprived enzyme−carbon nanotube hybrids incationic reverse micelles. Langmuir 25:4421–8.

265. Shin, J.Y, Premkumar, T., and Geckeler, K.E. 2008.Dispersion of single-walled carbon nanotubes by usingsurfactants: Are the type and concentration important. ChemEur J 14:6044–8.

266. Doe, C., Choi, S.M., Kline, S.R., Jang, H.S., and Kim,T.H. 2008. Charged rod-like nanoparticles assistingsingle-walled carbon nanotube dispersion in water. AdvFunct Mater 18:2685–91.

267. Wang, Q., Han, Y., Wang, Y., Qin, Y., and Guo, Z.X.2008. Effect of surfactant structure on the stability ofcarbon nanotubes in aqueous solution. J Phy Chem B112:7227–33.

268. Park, I., Lee, W., and Kim, J. 2007. Selectivesequestering of multi-walled carbon nanotubes inselfassembled block copolymer. Sensor Act B Chem 126:301–5.

269. Duesberg, G.S., Burghard, M., Muster, J., and Philipp,G. 1998. Separation of carbon nanotubes by size exclusionchromatography. Chem Commun (3):435–6.

270. Kim, Y., Hong, S., Jung, S., Strano, M.S., Choi, J.,and Baik, S. 2006. Dielectrophoresis of surfaceconductance modulated single-walled carbon nanotubes usingcatanionic surfactants. J Phys Chem B 110:1541–5.

271. Kim, S.K., Park, W., Kang, M., and Jin, H.J. 2008.Multiple light scattering measurement and stabilityanalysis of aqueous carbon nanotube dispersions. J Phy ChemSolids 69:1209–12.

272. Lisunova, M.O., Lebovka, N.I., Olexander, V. M., andBoiko, Y.P. 2006. Stability of the aqueous suspensions ofnanotubes in the presence of nonionic surfactant. J ColloidInterface Sci 299:740–6.

273. Kim, S.S., Lee, D.L., Shah, J., and Pinnavaia, T.J.2003. Nanocasting of carbon nanotubes: In-situgraphitization of a low-cost mesostructured silicatemplated by non-ionic surfactant micelles. Chem Commun

(12):1436–7.

274. Chen, R.J., Zhang, Y., Wang, D., and Dai, H. 2001.Noncovalent sidewall functionalization of singlewalledcarbon nanotubes for protein immobilization. J Am Chem Soc123:3838–9.

275. Guo, Z., Sadler, P.J., and Tsang, S.C. 1998.Immobilization and visualization of DNA and proteins oncarbon nanotubes. Adv Mater 10:701–3.

276. Barone, P.W. and Strano, M.S. 2006. Reversible controlof carbon nanotube aggregation for a glucose af�nitysensor. Angew Chem Int Ed 45:8138–41.

277. Munenori, N., Sugikawa, K., Kaneko, K., and Shinkai,S. 2008. Creation of hierarchical carbon nanotubeassemblies through alternative packing of complementarysemi-arti�cial β-1,3-glucan/carbon nanotube composites.Chem Eur J 14:2398–404.

278. Goodwin, A.P., Tabakman, S.M., Welsher, K., Sherlock,S.P., Prencipe, G., and Dai, H. 2009. Phospholipid−dextranwith a single coupling point: A useful amphiphile forfunctionalization of nanomaterials. J Am Chem Soc131:289–96.

279. Rouse, J.H. 2005. Polymer-assisted dispersion ofsingle-walled carbon nanotubes in alcohols andapplicability toward carbon nanotube/sol−gel composite.Langmuir 21:1055–61.

280. Wang, D., Ji, W.X., Li, Z.C., and Chen. L. 2006. Abiomimetic “polysoap” for single-walled carbon nanotubedispersion. J Am Chem Soc 128:6556–7.

281. Yuan, W.Z., Sun, J.Z., and Dong, Y. 2006. Wrappingcarbon nanotubes in pyrene-containingpoly(phenylacetylene) chains: Solubility, stability, lightemission, and surface photovoltaic properties.Macromolecules 39:8011–20.

282. Steuerman, D.W., Alexander, S., Narizzano, R. et al.2002. Interactions between conjugated polymers andsingle-walled carbon nanotubes. J Phys Chem B 106:3124–30.For cataloging purposes ly

283. Coleman, K.S., Bailey, S.R., Fogden, S., and Green,M.L.H. 2003. Functionalization of single-walled carbonnanotubes via the bingel. J Am Chem Soc 125:8722–3.

284. Touhara, H., Inahara, J., Mizuno, T. et al. 2002.Property control of new forms of carbon materials by�uorination. J Fluor Chem 114:181.

285. Marcoux, P.R., Schreiber, J., Batail, P. et al. 2002.A spectroscopic study of the �uorination and de�uorinationreactions on single-walled carbon nanotubes. Phys Chem ChemPhys 4:2278–85.

286. Kónya, Z., Vesselenyi, I., Niesz, K. et al. 2002.Large scale production of short functionalized carbonnanotubes Chem Phys Lett 360:429.

287. Hou, P.X., Bai, S., Yang, Q.H., Liu, C., and Cheng,H.M. 2002. Multi-step puri�cation of carbon nanotubes.Carbon 40:81–5.

288. Chang, J.W., Wu, H.Y., Hwang, G.L., and Su, T.Y. 2008.Ultrasonication-assisted surface functionalization ofdouble-walled carbon nanotubes with azobis-type radicalinitiators. J Mater Chem 18:3972–6.

289. Peng, H., Paul, R., Khabashesku, V.N., and Margrave,J.L. 2003. Sidewall functionalization of singlewalledcarbon nanotubes with organic peroxides. Chem Commun(3):362–3.

290. Jana, R.N. and Cho, J.W. 2008. Thermal stability,crystallization behavior, and phase morphology ofpoly(ε-caprolactone)diol-grafted-multiwalled carbonnanotubes. J Appl Polym Sci 110:1550–8.

11 Chapter 11 - Emerging Technology inAerospace Engineering: Polymer-BasedSelf-Healing Materials

Amamoto, Y., J. Kamada, H. Otsuka, A. Takahara, and K.Matyjaszewski. 2011. Repeatable photoinduced selfhealing ofcovalently cross-linked polymers through reshuf�ing oftrithiocarbonate units. Angewandte Chemie InternationalEdition 50(7):1660–1663.

Andreeva, D. V., D. Fix, H. Möhwald, and D. G. Shchukin.2008. Self-healing anticorrosion coatings based onpH-sensitive polyelectrolyte/inhibitor sandwichlikenanostructures. Advanced Materials 20(14):2789–2794.

Anna, C. B. 2007. Modeling self-healing materials.Materials Today 10(9):18–23.

Arshady, R. 1989. Preparation of microspheres andmicrocapsules by interfacial polycondensation techniques.Journal of Microencapsulation 6(1):13–28.

Avramova, N. 1993. Study of the healing-process of polymerswith different chemical-structure and chain mobility.Polymer 34(9):1904–1907.

Baker, A., S. Dutton, and D. Kelly. 2004. CompositeMaterials for Aircraft Structures. Second Edition,American Institute of Aeronautics and Astronautics, Reston,VA.

Balazs, A.C., T. Emrick, and T. P. Russell. 2006.Nanoparticle polymer composites: Where two small worldsmeet. Science 314(5802):1107–1110.

Beiermann, B. A., M. W. Keller, and N. R. Sottos. 2009.Self-healing �exible laminates for resealing of puncturedamage. Smart Materials and Structures 18(8):085001.

Blaiszik, B. J., M. Baginska, S. R. White, and N. R.Sottos. 2010a. Autonomic recovery of �ber/matrixinterfacial bond strength in a model composite. AdvancedFunctional Materials 20(20):3547–3554.

Blaiszik, B. J., M. M. Caruso, D. A. McIlroy, J. S. Moore,S. R. White, and N. R. Sottos. 2009. Microcapsules �lledwith reactive solutions for self-healing materials. Polymer50(4):990–997.

Blaiszik, B. J., S. L. B. Kramer, S. C. Olugebefola, J. S.

Moore, N. R. Sottos, and S. R. White. 2010b. Selfhealingpolymers and composites. Annual Review of MaterialsResearch 40:179–211.

Bleay, S. M., C. B. Loader, V. J. Hawyes, L. Humberstone,and P. T. Curtis. 2001. A smart repair system for polymermatrix composites. Composites Part—Applied Science andManufacturing 32(12):1767–1776.

Bonatz, E., K. Dietrich, H. Herma, R. Nastke, M. Walter,and W. Teige. 1989. Amino resin microcapsules. 3. Releaseproperties. Acta Polymerica 40(11):683–690.

Bond, I. P., R. S. Trask, and H. R. Williams. 2008.Self-healing �ber-reinforced polymer composites. MrsBulletin 33(8):770–774.

Brandon, E. J., M. Vozoff, E. A. Kolawa, et al. 2011.Structural health management technologies for in�atable/deployable structures: Integrating sensing andself-healing. Acta Astronautica 68(7–8):883–903.

Brochu, A. B. W., S. L. Craig, and W. M. Reichert. 2011.Self-healing biomaterials. Journal of Biomedical MaterialsResearch Part A 96A (2):492–506.

Brown, M. 2011. Self-healing ‘magic skin’, a condom foraircraft. Wired UK, 13 April.

Brown, E. N., M. R. Kessler, N. R. Sottos, and S. R. White.2003a. In situ poly(urea-formaldehyde) microencapsulationof dicyclopentadiene. Journal of Microencapsulation20(6):719–730.

Brown, E. N., J. S. Moore, S. R. White, and N. R. Sottos.2003b. Fracture and fatigue behavior of a self-healingpolymer composite. Material Research Society, 735(1):101–106.

Brown, E. N., N. R. Sottos, and S. R. White. 2002. Fracturetesting of a self-healing polymer composite. ExperimentalMechanics 42(4):372–379.

Brown, E. N., S. R. White, and N. R. Sottos. 2004.Microcapsule induced toughening in a self-healing polymercomposite. Journal of Materials Science 39(5):1703–1710.

Brown, E. N., S. R. White, and N. R. Sottos. 2005a.Retardation and repair of fatigue cracks in a microcapsuletoughened epoxy composite—Part 1: Manual in�ltration.

Composites Science and Technology 65(15–16):2466–2473.

Brown, E. N., S. R. White, and N. R. Sottos. 2005b.Retardation and repair of fatigue cracks in a microcapsuletoughened epoxy composite—Part II: In situ self-healing.Composites Science and Technology 65(15–16):2474–2480.

Brown, E.N., S. R. White, and N. R. Sottos. 2006. Fatiguecrack propagation in microcapsule-toughened epoxy. Journalof Materials Science 41(19):6266–6273.

Burattini, S., H. M. Colquhoun, J. D. Fox, et al. 2009a. Aself-repairing, supramolecular polymer system: Healabilityas a consequence of donor-acceptor pi-pi stackinginteractions. Chemical Communications (44):6717–6719.

Burattini, S., H. M. Colquhoun, B. W. Greenland, and W.Hayes. 2009b. A novel self-healing supramolecular polymersystem. Faraday Discussions 143:251–264.

Burattini, S., B. W. Greenland, D. Chappell, H. M.Colquhoun, and W. Hayes. 2010a. Healable polymericmaterials: A tutorial review. Chemical Society Reviews39(6):1973–1985.

Burattini, S., B. W. Greenland, W. Hayes, M. E. Mackay, S.J. Rowan, and H. M. Colquhoun. 2011. A Supramolecularpolymer based on tweezer-type π−π stacking interactions:Molecular design for healability and enhanced toughness.Chemistry of Materials 23(1):6–8.

Burattini, S., B. W. Greenland, D. H. Merino, et al. 2010b.A healable supramolecular polymer blend based on aromaticπ−π stacking and hydrogen-bonding interactions. Journal ofthe American Chemical Society 132(34):12051–12058.

Burnworth, M., L. M. Tang, J. R. Kumpfer, et al. 2011.Optically healable supramolecular polymers. Nature472(7343):334–337.

Cakhshaee, M., R. A. Pethrick, H. Pashid, and D. C.Sherrington. 1985. Microencapsulation techniques, factorsin�uencing encapsulation ef�ciency. Polymer Commun.26:185–192.

Cardenas-Daw, C., A. Kroeger, W. Schaertl, P. Froimowicz,and K. Landfester. 2011. Reversible photocycloadditions, apowerful tool for tailoring (Nano) materials.Macromolecular Chemistry and Physics 213(2):144–156.

Carlson, J. A., J. M. English, and D. J. Coe. 2006. A�exible, self-healing sensor skin. Smart Materials &Structures 15(5):N129–N135.

Caruso, M. M., B. J. Blaiszik, S. R. White, N. R. Sottos,and J. S. Moore. 2008. Full recovery of fracture toughnessusing a nontoxic solvent-based self-healing system.Advanced Functional Materials 18(13):1898–1904.

Caruso, M. M., D. A. Davis, Q. Shen, et al. 2009.Mechanically-induced chemical changes in polymericmaterials. Chemical Reviews 109(11):5755–5798. For ctaloging u poses only

Caruso, M. M., D. A. Delafuente, V. Ho, N. R. Sottos, J. S.Moore, and S. R. White. 2007. Solvent-promotedself-healing epoxy materials. Macromolecules40(25):8830–8832.

Chen, X. 2003. Novel polymers with thermally controlledcovalent cross-linking. PhD Thesis, University ofCalifornia, Los Angeles.

Chen, X. X., M. A. Dam, K. Ono, et al. 2002. A thermallyre-mendable cross-linked polymeric material. Science295(5560):1698–1702.

Chen, X., F. Wudl, A. K. Mal, H. Shen, and S. R. Nutt.2003. New thermally remendable highly cross-linkedpolymeric materials. Macromolecules 36(6):1802–1807.

Cho, S. H., H. M. Andersson, S. R. White, N. R. Sottos, andP. V. Braun. 2006. Polydimethylsiloxane-based self-healingmaterials. Advanced Materials 18(8):997.

Cho, S. H., S. R. White, and P. V. Braun. 2009.Self-healing polymer coatings. Advanced Materials21(6):645–649.

Chung, C. M., Y. S. Roh, S. Y. Cho, and J. G. Kim. 2004.Crack healing in polymeric materials via photochemical 2+2cycloaddition. Chemistry of Materials 16(21):3982–3984.

Coates, G. W., P. D. Hustad, and S. Reinartz. 2002.Catalysts for the living insertion polymerization ofalkenes: Access to new polyole�n architectures usingZiegler-Natta chemistry. Angewandte Chemie-InternationalEdition 41(13):2236–2257.

Cordier, P., F. Tournilhac, C. Soulie-Ziakovic, and L.

Leibler. 2008. Self-healing and thermoreversible rubberfrom supramolecular assembly. Nature 451(7181):977–980.

Cosco, S., V. Ambrogi, P. Musto, and C. Carfagna. 2007.Properties of poly(urea-formaldheyde) microcapsulescontaining an epoxy resin. Journal of Applied PolymerScience 105(3):1400–1411.

De Gennes, P. G. 1971. Reptation of a polymer chain in thepresence of �xed obstacles. The Journal of ChemicalPhysics 55(2):572–579.

Dietrich, K., E. Bonatz, H. Geistlinger, et al. 1989a.Amino resin microcapsules.2. Preparation and morphology.Acta Polymerica 40(5):325–331.

Dietrich, K., E. Bonatz, R. Nastke, H. Herma, M. Walter,and W. Teige. 1990. Amino resin microcapsules.4.Surface-tension of the resins and mechanism of capsulefromation. Acta Polymerica 41(2):91–95.

Dietrich, K., H. Herma, R. Nastke, E. Bonatz, and W. Teige.1989b. Amino resin microcapsules.1. Literature and patentreview. Acta Polymerica 40(4):243–250.

Doi, M., and S. F. Edwards. 1978. Dynamics of concentratedpolymer systems. Part 1—Brownian motion in the equilibriumstate. Journal of the Chemical Society, FaradayTransactions 2: Molecular and Chemical Physics74:1789–1801.

Dorworth, L. C., G. L. Gardiner, and G. M. Mellema. 2010.Essentials of Advanced Composite Fabrication and Repair:Aviation Supplies & Academics, Inc., Newcastle, WA.

Dry, C. 1996. Procedures developed for self-repair ofpolymer matrix composite materials. Composite Structures35(3):263–269.

Dry, C. M. 2006. Self-repairing, reinforced matrixmaterials (Individual U) US:7022179-B1.

Eisenberg, A., and J. S. Kim. 1998. Introduction toIonomers: Wiley-Interscience, New York, NY.

Eisenberg, A., and M. Rinaudo. 1990. Polyelectrolytes andionomers. Polymer Bulletin 24(6):671–671.

Fall, R. 2001. Puncture Reversal of EthyleneIonomers-Mechanistic Studies. Virginia Polytechnic

Institute and State University, Blacksburg.

Fischer, H.. 2010. Self-repairing material systems—A dreamor a reality? Natural Science 2(8):873–901.

García, S. J., H. R. Fischer, P. A. White, et al. 2011.Self-healing anticorrosive organic coating based on anencapsulated water reactive silyl ester: Synthesis andproof of concept. Progress in Organic Coatings70(2–3):142–149.

García, S. J., M. T. Rodríguez, R. Izquierdo, and J. Suay.2007. Evaluation of cure temperature effects incataphoretic automotive primers by electrochemicaltechniques. Progress in Organic Coatings 60(4):303–311.

Ghosh, S. K. 2006. Functional Coatings: by PolymerMicroencapsulation: Wiley-VCH, Weinheim, Germany.

Ghosh, S. K. 2009. Self-Healing Materials: Fundamentals,Design Strategies, and Applications: Wiley-VCH, Weinheim,Germany.

Ghosh, B., K. V. Chellappan, and M. W. Urban. 2011.Self-healing inside a scratch of oxetane-substitutedchitosan-polyurethane (OXE-CHI-PUR) networks. Journal ofMaterials Chemistry 21(38):14473–14486.

Ghosh, B., and M. W. Urban. 2009. Self-repairingoxetane-substituted chitosan polyurethane networks. Science323(5920):1458–1460.

Goethals, E. J., and F. Du Prez. 2007. Carbocationicpolymerizations. Progress in Polymer Science 32(2):220–246.

Gupta, S., Q. L. Zhang, T. Emrick, A. C. Balazs, and T. P.Russell. 2006. Entropy-driven segregation of nanoparticlesto cracks in multilayered composite polymer structures.Nature Materials 5(3):229–233.

Hager, M. D., P. Greil, C. Leyens, S. van der Zwaag, and U.S. Schubert. 2010. Self-healing materials. AdvancedMaterials 22:5424–5430.

Hamilton, A. R., N. R. Sottos, and S. R. White. 2010.Self-healing of internal damage in synthetic vascularmaterials. Advanced Materials 22(45):5159–5163.

Hamilton, A. R., N. R. Sottos, and S. R. White. 2011.Pressurized vascular systems for self-healing materials.

Journal of the Royal Society Interface. doi:10.1098/rsif.2011.0508.

Hansen, C. J., W. Wu, K. S. Toohey, N. R. Sottos, S. R.White, and J. A. Lewis. 2009. Self-healing materials withinterpenetrating microvascular networks. Advanced Materials21(41):4143–4147.

Hayes, S. A., F. R. Jones, K. Marshiya, and W. Zhang.2007a. A self-healing thermosetting composite material.Composites Part A—Applied Science and Manufacturing38(4):1116–1120.

Hayes, S. A., W. Zhang, M. Branthwaite, and F. R. Jones.2007b. Self-healing of damage in �bre-reinforcedpolymer-matrix composites. Journal of the Royal SocietyInterface 4(13):381–387.

Hsieh, H. C., T. J. Yang, and S. Lee. 2001. Crack healingin poly(methyl methacrylate) induced by co-solvent ofmethanol and ethanol. Polymer 42(3):1227–1241.

Huang, C. Y., R. S. Trask, and I. P. Bond. 2010.Characterization and analysis of carbon �bre-reinforcedpolymer composite laminates with embedded circularvasculature. Journal of the Royal Society Interface7(49):1229–1241.

Huang, M., and J. Yang. 2011. Facile microencapsulation ofHDI for self-healing anticorrosion coatings. Journal ofMaterials Chemistry 21(30):11123–11130.

Hucker, M., I. Bond, S. Bleay, and S. Haq. 2003a.Experimental evaluation of unidirectional hollow glass�bre/ epoxy composites under compressive loading.Composites Part A—Applied Science and Manufacturing34(10):927–932.

Hucker, M., I. Bond, S. Bleay, and S. Haq. 2003b.Investigation into the behaviour of hollow glass �brebundles under compressive loading. Composites PartA—Applied Science and Manufacturing 34(11):1045–1052.

Hurley, D. A., and D. R. Huston. 2011. Coordinated sensingand active repair for self-healing. Smart Materials andStructures 20(2):025010.

Ignác, C. 2005. Nature and properties of ionomerassemblies. II. Advances in Colloid and Interface Science118(1–3):73–112.

Imaizumi, K., T. Ohba, Y. Ikeda, and K. Takeda. 2001.Self-repairing mechanism of polymer composite. MaterialsScience Research International 7(4):249–253.

Isaacs, H. S. 1988. The measurement of the galvaniccorrosion of soldered copper using the scanning vibratingelectrode technique. Corrosion Science 28(6):547–558.

Jackson, A. C., J. A. Bartelt, and P. V. Braun. 2011.Transparent self-healing polymers based on encapsulatedplasticizers in a thermoplastic matrix. Advanced FunctionalMaterials 21(24):4705–4711.

Jin, H., C. L. Mangun, D. S. Stradley, J. S. Moore, N. R.Sottos, and S. R. White. 2012. Self-healing thermosetusing encapsulated epoxy-amine healing chemistry. Polymer53(2):581–587.

Jin, H., C. L. Mangun, D. S. Stradley, N. R. Sottos, and S.R. White. 2011b. Epoxy-amine based self-healing in hightemperature cured epoxy. In The 3rd InternationalConference on Self-Healing Materials. Bath, UK.

Jin, H., G. M. Miller, N. R. Sottos, and S. R. White.2011a. Fracture and fatigue response of a self-healingepoxy adhesive. Polymer 52(7):1628–1634.

Johnson, S. B., T. Gormley, S. Kessler, et al. 2011. SystemHealth Management: With Aerospace Applications: John Wiley& Sons, Hoboken, NJ.

Jones, A. S., J. D. Rule, J. S. Moore, N. R. Sottos, and S.R. White. 2007. Life extension of self-healing polymerswith rapidly growing fatigue cracks. Journal of the RoyalSociety Interface 4(13):395–403.

Jud, K., and H. H. Kausch. 1979. Load-transfer throughchain molecules after interpenetration at interfaces.Polymer Bulletin 1(10):697–707.

Jud, K., H. H. Kausch, and J. G. Williams. 1981.Fracture-mechanics studies of crack healing and welding ofpolymers. Journal of Materials Science 16(1):204–210.

Kalista, S. J. 2003. Self-healing of thermoplasticpoly(ethylene-co-methacrylic acid) copolymers followingprojectile puncture. Master Thesis, Virginia PolytechnicInstitute and State University, Blacksburg.

Kalista, S. J. 2007. Self-healing ofpoly(ethylene-co-methacrylic acid) copolymers followingprojectile puncture. Mechanics of Advanced Materials andStructures 14(5):391–397.

Kamphaus, J. M., J. D. Rule, J. S. Moore, N. R. Sottos, andS. R. White. 2008. A new self-healing epoxy with tungsten(VI) chloride catalyst. Journal of the Royal SocietyInterface 5(18):95–103.

Kan, M., C.-s., Hiroshi Hyodo, and Setagayapku. 1969.Pressure sentitive copying sheet and the productionthereof. United States Patent 3432327.

Kavitha, A. A., and N. K. Singha. 2009. “Click chemistry”in tailor-made polymethacrylates bearing reactive furfurylfunctionality: A new class of self-healing polymericmaterial. ACS Applied Materials & Interfaces1(7):1427–1436. For aloging pu p es o l

Kawagoe, M., M. Nakanishi, J. Qiu, and M. Morita. 1997.Growth and healing of a surface crack in poly(methylmethacrylate) under case II diffusion of methanol. Polymer38(24):5969–5975.

Keller, M. W. 2007. A self-healing poly(dimethyl siloxane)elastomer. PhD Thesis, University of Illinois atUrbana-Champaign, Urbana.

Keller, M. W., S. R. White, and N. R. Sottos. 2007. Aself-healing poly(dimethyl siloxane) elastomer. AdvancedFunctional Materials 17(14):2399–2404.

Keller, M. W., S. R. White, and N. R. Sottos. 2008. Torsionfatigue response of self-healing poly(dimethylsiloxane)elastomers. Polymer 49(13–14):3136–3145.

Kessler, M. R. 2002. Chacterization and Performance of aSelf-Healing Composite Material. University of Illinois atUrbana-Chanpaign, Urbana.

Kessler, M. R. 2007. Self-healing: A new paradigm inmaterials design. Proceedings of the Institution ofMechanical Engineers Part G-Journal of AerospaceEngineering 221(G4):479–495.

Kessler, M. R., N. R. Sottos, and S. R. White. 2003.Self-healing structural composite materials. CompositesPart A: Applied Science and Manufacturing 34(8):743–753.

Kessler, M. R., and S. R. White. 2001. Self-activatedhealing of delamination damage in woven composites.Composites Part A—Applied Science and Manufacturing32(5):683–699.

Kessler, M. R., and S. R. White. 2002. Cure kinetics of thering-opening metathesis polymerization ofdicyclopentadiene. Journal of Polymer Science, Part A:Polymer Chemistry 40(14):2373–2383.

Khramov, A. N., N. N. Voevodin, V. N. Balbyshev, and M. S.Donley. 2004. Hybrid organo-ceramic corrosion protectioncoatings with encapsulated organic corrosion inhibitors.Thin Solid Films 447–448:549–557.

Kim, H. J., Ki-Jun Lee, and Hong, H. Lee. 1996. Healing offractured polymers by interdiffusion. Polymer37(20):4593–4597.

Kim, Y. H., and R. P. Wool. 1983. A theory of healing at apolymer-polymer interface. Macromolecules 16(7):1115–1120.

Kirkby, E. L., J. D. Rule, V. L. Michaud, N. R. Sottos, S.R. White, and J. A. E. Manson. 2008. Embedded shape-memoryalloy wires for improved performance of self-healingpolymers. Advanced Functional Materials 18(15):2253–2260.

Kolmakov, G.V., R. Revanur, R. Tangirala, et al. 2010.Using nanoparticle-�lled microcapsules for site-speci�chealing of damaged substrates: Creating a “repair-and-go”system. ACS Nano 4(2):1115–1123.

Kroes, M. J., F. Delp, and W. A. Watkins. 2002. AircraftMaintenance and Repair: McGraw-Hill Education, New York,NY.

Kuang, F, T. Shi, J. Wang, and F. Jia. 2009.Microencapsulation technology for thiourea corrosioninhibitor. Journal of Solid State Electrochemistry13(11):1729–1735.

Kumar, A., L. D. Stephenson, and J. N. Murray. 2006.Self-healing coatings for steel. Progress in OrganicCoatings 55(3):244–253.

Kwok, N., and H. T. Hahn. 2007. Resistance heating forself-healing composites. Journal of Composite Materials41(13):1635–1654.

Lee, J. Y., G. A. Buxton, and A. C. Balazs. 2004. Using

nanoparticles to create self-healing composites. Journalof Chemical Physics 121(11):5531–5540.

Lee, J.-Y., Q. Zhang, T. Emrick, and A. J. Crosby. 2006.Nanoparticle alignment and repulsion during failure ofglassy polymer nanocomposites. Macromolecules39(21):7392–7396.

Lewis, J. A. 2006. Direct ink writing of 3D functionalmaterials. Advanced Functional Materials 16(17): 2193–2204.

Lin, C. B., S. B. Lee, and K. S. Liu. 1990.Methanol-induced crack healing in poly(methylmethacrylate). Polymer Engineering and Science30(21):1399–1406.

Ling, J., M. Z. Rong, and M. Q. Zhang. 2011. Coumarinimparts repeated photochemical remendability topolyurethane. Journal of Materials Chemistry21(45):18373–18380.

Liu, Y. L., and Y. W. Chen. 2007. Thermally reversiblecross-linked polyamides with high toughness andself-repairing ability from maleimide- andfuran-functionalized aromatic polyamides. MacromolecularChemistry and Physics 208(2):224–232.

Liu, Y. L., and C. Y. Hsieh. 2006. Crosslinked epoxymaterials exhibiting thermal remendablility andremovability from multifunctional maleimide and furancompounds. Journal of Polymer Science Part A— PolymerChemistry 44(2):905–913.

Liu, C. K., T. J. Yang, J. S. Shen, and S. Lee. 2008. Somerecent results on crack healing of poly(methylmethacrylate). Engineering Fracture Mechanics75(17):4876–4885.

Liu, X., X. Sheng, J. Keun Lee, and M. R. Kessler. 2009.Synthesis and characterization of melamine-ureaformaldehydemicrocapsules containing ENB-based self-healing agents.Macromolecular Materials and Engineering 294(6–7):389–395.

Loveday, D., P. Peterson, and B.R.G. Instruments. 2004.Evaluation of organic coatings with electrochemicalimpedance spectroscopy. Journal of Coatings Technology8:46–52.

Luo, X. F., R. Q. Ou, D. E. Eberly, A. Singhal, W.Viratyaporn, and P. T. Mather. 2009. A

thermoplastic/thermoset blend exhibiting thermal mendingand reversible adhesion. ACS Applied Materials & Interfaces1(3):612–620.

Marieb, E. N. 2009. Essentials of Human Anatomy &Physiology. 9th ed: Pearson/Benjamin Cummings, New York,NY.

Mauldin, T. C., and M. R. Kessler. 2010. Self-healingpolymers and composites. International Materials Reviews55(6):317–346.

McIlroy, D. A., B. J. Blaiszik, M. M. Caruso, S. R. White,J. S. Moore, and N. R. Sottos. 2010. Microencapsulation ofa reactive liquid-phase amine for self-healing epoxycomposites. Macromolecules 43(4):1855–1859.

Meng, L. M., Y. Chao Yuan, M. Z. Rong, and M. Q. Zhang.2010. A dual mechanism single-component selfhealingstrategy for polymers. Journal of Materials Chemistry20(29):6030–6038.

Meure, S., S. Furman, and S. Khor. 2010a. Polyethylene-co-(methacrylic acid) healing agents for mendablecarbon �ber laminates. Macromolecular Materials andEngineering 295(5):420–424.

Meure, S., D. Y. Wu, and S. Furman. 2009.Polyethylene-co-methacrylic acid healing agents formendable epoxy resins. Acta Materialia 57(14):4312–4320.

Meure, S., D. Y. Wu, and S. A. Furman. 2010b. FTIR study ofbonding between a thermoplastic healing agent and amendable epoxy resin. Vibrational Spectroscopy 52(1):10–15.

Miyano, Y., and M. Nakada. 2011. In Advanced AcceleratedTesting Methodology for Life Prediction of CFRP laminatesTime Dependent Constitutive Behavior and Fracture/FailureProcesses, Volume 3, eds. by T. Proulx: Springer, NewYork.

Miyano, Y., M. Nakada, and H. Cai. 2008. Formulation oflong-term creep and fatigue strengths of polymer compositesbased on accelerated testing methodology. Journal ofComposite Materials 42(18):1897–1919.

Moll, J. L. 2011. Self-sealing of thermal fatigue andmechanical damage in �ber-reinforced composite materials.PhD Thesis, University of Illinois at Urbana-Champaign,Urbana.

Moll, J. L., S. R. White, and N. R. Sottos. 2010. Aself-sealing �ber-reinforced composite. Journal ofComposite Materials 44(22):2573–2585.

Montarnal, D., P. Cordier, C. Soulie-Ziakovic, F.Tournilhac, and L. Leibler. 2008. Synthesis of self-healingsupramolecular rubbers from fatty acid derivatives,diethylenetriamine, and urea. Journal of Polymer SciencePart A—Polymer Chemistry 46(24):7925–7936.

Mostovoy, S., P. B. Crosley, and E. J. Ripling. 1967. Useof crack-line load specimens for measuring planestrainfracture toughness. Journal of Materials 2(3):661–681.

Motuku, M., U. K. Vaidya, and G. M. Janowski. 1999.Parametric studies on self-repairing approaches for resininfused composites subjected to low velocity impact. SmartMaterials & Structures 8(5):623–638.

Murphy, E. B., E. Bolanos, C. Schaffner-Hamann, F. Wudl, S.R. Nutt, and M. L. Auad. 2008. Synthesis andcharacterization of a single-component thermally remendablepolymer network: Staudinger and Stille revisited.Macromolecules 41(14):5203–5209.

Murphy, E. B., and F. Wudl. 2010. The world of smarthealable materials. Progress in Polymer Science35(1–2):223–251.

Nguyen, T. Q., H. H. Kausch, K. Jud, and M. Dettenmaier.1982. Crack-healing in crosslinked styrene-coacrylonitrile.Polymer 23(9):1305–1310.

Norris, C. J., I. P. Bond, and R. S. Trask. 2011a.Interactions between propagating cracks and bioinspiredselfhealing vascules embedded in glass �bre reinforcedcomposites. Composites Science and Technology71(6):847–853.

Norris, C. J., I. P. Bond, and R. S. Trask. 2011b. The roleof embedded bioinspired vasculature on damage formation inself-healing carbon �bre reinforced composites. CompositesPart A—Applied Science and Manufacturing 42(6):639–648.

Olugebefola, S. C., A. M. Aragon, C. J. Hansen, et al.2010. Polymer microvascular network composites. Journal ofComposite Materials 44(22):2587–2603.

Osswald, T., and G. Menges. 2003. Failure and damage of

polymers. In Materials Science of Polymers for Engineers,eds. by T. Osswald, and G. Menges: Hanser Publishers,Munich: pp. 447–519.

Pang, J. W. C., and I. P. Bond. 2005a. ‘Bleedingcomposites’—Damage detection and self-repair using abiomimetic approach. Composites Part A—Applied Science andManufacturing 36(2):183–188.

Pang, J. W. C., and I. P. Bond. 2005b. A hollow �brereinforced polymer composite encompassing self-healing andenhanced damage visibility. Composites Science andTechnology 65(11–12):1791–1799.

Paris, P. C., M. P. Gomez, and W. E. Anderson. 1961. TheTrend in Engineering at the University of Washington13:9–14. For catalogi g purp ses only

Park, J. S., T. Darlington, A. F. Starr, K. Takahashi, J.Riendeau, and H. T. Hahn. 2010. Multiple healing effect ofthermally activated self-healing composites based onDiels-Alder reaction. Composites Science and Technology70(15):2154–2159.

Park, J. S., K. Takahashi, Z. H. Guo, et al. 2008. Towardsdevelopment of a self-healing composite using a mendablepolymer and resistive heating. Journal of CompositeMaterials 42(26):2869–2881.

Park, J.-H., and P. V. Braun. 2010. Coaxial electrospinningof self-healing coatings. Advanced Materials22(4):496–499.

Patel, A. J., N. R. Sottos, E. D. Wetzel, and S. R. White.2010. Autonomic healing of low-velocity impact damage in�ber-reinforced composites. Composites Part A—AppliedScience and Manufacturing 41(3):360–368.

Peterson, A. M., R. E. Jensen, and G. R. Palmese. 2009.Reversibly cross-linked polymer gels as healing agents forepoxy-amine thermosets. ACS Applied Materials & Interfaces1(5):992–995.

Peterson, A. M., R. E. Jensen, and G. R. Palmese. 2010.Room-temperature healing of a thermosetting polymer usingthe diels-alder reaction. ACS Applied Materials &Interfaces 2(4):1141–1149.

Plaisted, T. A., and S. Nemat-Nasser. 2007. Quantitativeevaluation of fracture, healing and re-healing of a

reversibly cross-linked polymer. Acta Materialia55(17):5684–5696.

Rule, J. D., E. N. Brown, N. R. Sottos, S. R. White, and J.S. Moore. 2005. Wax-protected catalyst microspheres foref�cient self-healing materials. Advanced Materials17(2):205–208.

Rule, J. D., N. R. Sottos, and S. R. White. 2007. Effect ofmicrocapsule size on the performance of self-healingpolymers. Polymer 48(12):3520–3529.

Rybtchinski, B. 2011. Adaptive supramolecular nanomaterialsbased on strong noncovalent interactions. ACS Nano5(9):6791–6818.

Samadzadeh, M., S. Hatami Boura, M. Peikari, A. Ashra�, andM. Kasiriha. 2011. Tung oil: An autonomous repairing agentfor self-healing epoxy coatings. Progress in OrganicCoatings 70(4):383–387.

Sanada, K., I. Yasuda, and Y. Shindo. 2006. Transversetensile strength of unidirectional �bre-reinforced polymersand self-healing of interfacial debonding. Plastics Rubberand Composites 35(2):67–72.

Sanyal, A. 2010. Diels–alder cycloaddition-cycloreversion:A powerful combo in materials design. MacromolecularChemistry and Physics 211(13):1417–1425.

Sauvant-Moynot, V., S. Gonzalez, and J. Kittel. 2008.Self-healing coatings: An alternative route foranticorrosion protection. Progress in Organic Coatings63(3):307–315.

Shchukin, D. G., and H. Möhwald. 2007. Self-repairingcoatings containing active nanoreservoirs. Small3(6):926–943.

Shchukin, D. G., M. Zheludkevich, K. Yasakau, S. Lamaka,M. G. S. Ferreira, and H. Möhwald. 2006. Layer-by-layerassembled nanocontainers for self-healing corrosionprotection. Advanced Materials 18(13):1672–1678.

Sperling, L. H. 2006. Introduction to Physical PolymerScience: Wiley-Interscience, Hoboken, NJ.

Sperling, L. H., A. Klein, and M. Sambasivam. 1996. Amolecular approach to the healing and fracture at polymerinterfaces: Current status of the literature. Journal of

Polymer Materials 13(1):1–13.

Staszewski, W. J., C. Boller, and G. R. Tomlinson. 2004.Health Monitoring of Aerospace Structures. Munich: JohnWiley & Sons, Ltd.

Suryanarayana, C., K. Chowdoji Rao, and D. Kumar. 2008.Preparation and characterization of microcapsulescontaining linseed oil and its use in self-healingcoatings. Progress in Organic Coatings 63(1):72–78.

Syrett, J. A., C. R. Becer, and D. M. Haddleton. 2010.Self-healing and self-mendable polymers. Polymer Chemistry1(7):978–987.

Szwarc, M. 1956. Living polymers. Nature178(4543):1168–1169.

Takeda, K., M. Tanahashi, and H. Unno. 2003. Self-repairingmechanism of plastics. Science and Technology of AdvancedMaterials 4(5):435–444.

Takeda, K., H. Unno, and M. Zhang. 2004. Polymer reactionin polycarbonate with Na 2 CO 3 . Journal of AppliedPolymer Science 93(2):920–926.

Teoh, S. H., H. Y. Chia, M. S. Lee, et al. 2010. Selfhealing composite for aircraft’s structural application.International Journal of Modern Physics B 24(1–2):157–163.

Therriault, D. 2003. Directed assembly of three-dimensionalmicrovascular networks. PhD Thesis, University of Illinoisat Urbana-Champaign, Urbana.

Therriault, D., R. F. Shepherd, S. R. White, and J. A.Lewis. 2005. Fugitive inks for direct-write assembly ofthree-dimensional microvascular networks. AdvancedMaterials 17(4):395–399.

Therriault, D., S. R. White, and J. A. Lewis. 2003. Chaoticmixing in three-dimensional microvascular networksfabricated by direct-write assembly. Nature Materials2(4):265–271.

Tian, Q., M. Z. Rong, M. Q. Zhang, and Y. C. Yuan. 2010.Optimization of thermal remendability of epoxy viablending. Polymer 51(8):1779–1785.

Tian, Q., Y. C. Yuan, M. Z. Rong, and M. Q. Zhang. 2009. Athermally remendable epoxy resin. Journal of Materials

Chemistry 19(9):1289–1296.

Tong, X.-M., T. Zhang, M.-Z. Yang, and Q. Zhang. 2010.Preparation and characterization of novel melamine modi�edpoly(urea–formaldehyde) self-repairing microcapsules.Colloids and Surfaces A: Physicochemical and EngineeringAspects 371(1–3):91–97.

Toohey, K. S., C. J. Hansen, J. A. Lewis, S. R. White, andN. R. Sottos. 2009a. Delivery of TwoPart Self-HealingChemistry via microvascular networks. Advanced FunctionalMaterials 19(9):1399–1405.

Toohey, K. S., N. R. Sottos, J. A. Lewis, J. S. Moore, andS. R. White. 2007. Self-healing materials withmicrovascular networks. Nature Materials 6(8):581–585.

Toohey, K. S., N. R. Sottos, and S. R. White. 2009b.Characterization of microvascular-based self-healingcoatings. Experimental Mechanics 49(5):707–717.

Trask, R. S., and I. P. Bond. 2006. Biomimetic self-healingof advanced composite structures using hollow glass �bres.Smart Materials & Structures 15(3):704–710.

Trask, R. S., and I. P. Bond. 2010. Bioinspired engineeringstudy of Plantae vascules for self-healing compositestructures. Journal of the Royal Society Interface7(47):921–931.

Trask, R. S., G. J. Williams, and I. P. Bond. 2007.Bioinspired self-healing of advanced composite structuresusing hollow glass �bres. Journal of the Royal SocietyInterface 4(13):363–371.

Tyagi, S., J. Y. Lee, G. A. Buxton, and A. C. Balazs. 2004.Using nanocomposite coatings to heal surface defects.Macromolecules 37(24):9160–9168.

Urban, M. W. 2009. Strati�cation, stimuli-responsiveness,self-healing, and signaling in polymer networks. Progressin Polymer Science 34(8):679–687.

Urrehman, A., and P. F. Thomason. 1993. The effect ofarti�cial fatigue-crack closure on fatigue-crack growth.Fatigue & Fracture of Engineering Materials & Structures16(10):1081–1090.

van der Zwaag, S., A. J. M. Schmets, and G. Zaken. 2007.Self Healing Materials: An Alternative Approach to 20

Centuries of Materials Science: Springer, Dordrecht,Netherlands.

Varley, R. J., S. Shen, and S. van der Zwaag. 2010. Theeffect of cluster plasticisation on the self healingbehaviour of ionomers. Polymer 51(3):679–686.

Varley, R. J., and S. van der Zwaag. 2008a. Towards anunderstanding of thermally activated self-healing of anionomer system during ballistic penetration. ActaMaterialia 56(19):5737–5750.

Varley, R. J., and S. van der Zwaag. 2008b. Development ofa quasi-static test method to investigate the origin ofself-healing in ionomers under ballistic conditions.Polymer Testing 27(1):11–19.

Wagg, D. 2007. Adaptive Structures: EngineeringApplications: John Wiley, Chichester, UK.

Wagner, M., and G. Norris. 2009. Boeing 787 Dreamliner:Zenith Press, Minneapolis, MN.

Wang, P. P., S. Lee, and J. P. Harmon. 1994.ethanol-induced crack healing in poly(methyl methacrylate).Journal of Polymer Science Part B—Polymer Physics32(7):1217–1227.

Wang, R. G., H. Y. Li, H. L. Hu, X. D. He, and W. B. Liu.2009. Preparation and characterization of self-healingmicrocapsules with poly(urea-formaldehyde) grafted epoxyfunctional group shell. Journal of Applied Polymer Science113(3):1501–1506.

Wang, H. P., Y. C. Yuan, M. Z. Rong, and M. Q. Zhang. 2010.Self-healing of thermoplastics via living polymerization.Macromolecules 43(2):595–598.

White, S. R., N. R. Sottos, P. H. Geubelle et al. 2001.Autonomic healing of polymer composites. Nature409(6822):794–797.

Williams, G. J., I. P. Bond, and R. S. Trask. 2009.Compression after impact assessment of self-healing CFRP.Composites Part A—Applied Science and Manufacturing40(9):1399–1406.

Williams, G., R. Trask, and I. Bond. 2007a. A self-healingcarbon �bre reinforced polymer for aerospace applications.Composites Part A—Applied Science and Manufacturing

38(6):1525–1532.

Williams, H. R., R. S. Trask, and I. P. Bond. 2007b.Self-healing composite sandwich structures. SmartMaterials & Structures 16(4):1198–1207.

Williams, H. R., R. S. Trask, and I. P. Bond. 2008a.Self-healing sandwich panels: Restoration of compressivestrength after impact. Composites Science and Technology68(15–16):3171–3177.

Williams, H. R., R. S. Trask, A. C. Knights, E. R.Williams, and I. P. Bond. 2008b. Biomimetic reliabilitystrategies for self-healing vascular networks inengineering materials. Journal of the Royal SocietyInterface 5(24):735–747.

Williams, H. R., R. S. Trask, P. M. Weaver, and I. P. Bond.2008c. Minimum mass vascular networks in multifunctionalmaterials. Journal of the Royal Society Interface5(18):55–65. For cataloging purpos only

Wilson, G. O., M. M. Caruso, N. T. Reimer, S. R. White, N.R. Sottos, and J. S. Moore. 2008a. Evaluation of rutheniumcatalysts for ring-opening metathesis polymerization-basedself-healing applications. Chemistry of Materials20(10):3288–3297.

Wilson, G. O., J. S. Moore, S. R. White, N. R. Sottos, andH. M. Andersson. 2008b. Autonomic healing of epoxy vinylesters via ring opening metathesis polymerization. AdvancedFunctional Materials 18(1):44–52.

Wool, R. P. 2008. Self-healing materials: A review. SoftMatter 4(3):400–418.

Wool, R. P., and K. M. O’Connor. 1981. A theory of crackhealing in polymers. Journal of Applied Physics52(10):5953–5963.

Wu, T., and S. Lee. 1994. Carbon tetrachloride-inducedcrack healing in polycarbonate. Journal of Polymer SciencePart B—Polymer Physics 32(12):2055–2064.

Wu, D. Y., S. Meure, and D. Solomon. 2008. Self-healingpolymeric materials: A review of recent developments.Progress in Polymer Science 33(5):479–522.

Xia, Y., and George M. Whitesides. 1998. Soft lithography.Annual Review of Materials Science 28(1):153–184.

Xiao, D. S., M. Z. Rong, and M. Q. Zhang. 2007. A novelmethod for preparing epoxy-containing microcapsules via UVirradiation-induced interfacial copolymerization inemulsions. Polymer 48(16):4765–4776.

Xiao, D. S., Y. C. Yuan, M. Z. Rong, and M. Q. Zhang.2009a. A facile strategy for preparing self-healingpolymer composites by incorporation of cationiccatalyst-loaded vegetable �bers. Advanced FunctionalMaterials 19(14):2289–2296.

Xiao, D. S., Y. C. Yuan, M. Z. Rong, and M. Q. Zhang.2009b. Hollow polymeric microcapsules: Preparation,characterization and application in holding borontri�uoride diethyl etherate. Polymer 50(2):560–568.

Xiao, D. S., Y. C. Yuan, M. Z. Rong, and M. Q. Zhang.2009c. Self-healing epoxy based on cationic chainpolymerization. Polymer 50(13):2967–2975.

Yamaguchi, M., S. Ono, and K. Okamoto. 2009. Interdiffusionof dangling chains in weak gel and its application toself-repairing material. Materials Science and EngineeringB—Advanced Functional Solid-State Materials162(3):189–194.

Yang, J. L., M. W. Keller, J. S. Moore, S. R. White, and N.R. Sottos. 2008. Microencapsulation of isocyanates forself-healing polymers. Macromolecules 41(24):9650–9655.

Yao, L., Y. C. Yuan, M. Z. Rong, and M. Q. Zhang. 2011a.Self-healing linear polymers based on RAFT polymerization.Polymer 52(14):3137–3145.

Yao, L., M. Z. Rong, M. Q. Zhang, and Y. C. Yuan. 2011b.Self-healing of thermoplastics via reversibleaddition-fragmentation chain transfer polymerization.Journal of Materials Chemistry 21(25):9060–9065.

Yin, T., M. Z. Rong, M. Q. Zhang, and G. C. Yang. 2007.Self-healing epoxy composites—Preparation and effect ofthe healant consisting of microencapsulated epoxy andlatent curing agent. Composites Science and Technology67(2):201–212.

Yin, T., M. Z. Rong, M. Q. Zhang, and J. Q. Zhao. 2009.Durability of self-healing woven glass fabric/epoxycomposites. Smart Materials & Structures18(7):074001–074007.

Yin, T., L. Zhou, M. Z. Rong, and M. Q. Zhang. 2008a.Self-healing woven glass fabric/epoxy composites with thehealant consisting of micro-encapsulated epoxy and latentcuring agent. Smart Materials & Structures17(1):015019–015026.

Yin, T., M. Zhi Rong, J. Wu, H. Chen, and M. Qiu Zhang.2008b. Healing of impact damage in woven glass fabricreinforced epoxy composites. Composites Part A: AppliedScience and Manufacturing 39(9):1479–1487.

Yoshie, N., M. Watanabe, H. Araki, and K. Ishida. 2010.Thermo-responsive mending of polymers crosslinked bythermally reversible covalent bond: Polymers frombisfuranic terminated poly(ethylene adipate) andtris-maleimide. Polymer Degradation and Stability95(5):826–829.

Yuan, L., A. J. Gu, and G. Z. Liang. 2008a. Preparation andproperties of poly(urea-formaldehyde) microcapsules �lledwith epoxy resins. Materials Chemistry and Physics110(2–3):417–425.

Yuan, L., G. Z. Liang, J. Q. Xie, L. Li, and J. Guo. 2006.Preparation and characterization of poly(urea-formaldehyde)microcapsules �lled with epoxy resins. Polymer47(15):5338–5349.

Yuan, Li, G.-Z. Liang, J.-Q. Xie, and S.-B. He. 2007.Synthesis and characterization of microencapsulateddicyclopentadiene with melamine–formaldehyde resins.Colloid & Polymer Science 285(7):781–791.

Yuan, Y. C., M. Z. Rong, and M. Q. Zhang. 2008b.Preparation and characterization of microencapsulatedpolythiol. Polymer 49(10):2531–2541.

Yuan, Y. C., M. Z. Rong, M. Q. Zhang, B. Chen, G. C. Yang,and X. M. Li. 2008c. Self-healing polymeric materialsusing epoxy/mercaptan as the healant. Macromolecules41(14):5197–5202.

Yuan, Y. C., M. Z. Rong, M. Q. Zhang, and G. C. Yang. 2009.Study of factors related to performance improvement ofself-healing epoxy based on dual encapsulated healant.Polymer 50(24):5771–5781.

Yuan, Y. C., T. Yin, M. Z. Rong, and M. Q. Zhang. 2008d.Self healing in polymers and polymer composites. Concepts,

realization and outlook: A review. Express Polymer Letters2(4):238–250.

Zako, M., and N. Takano. 1999. Intelligent material systemsusing epoxy particles to repair microcracks anddelamination damage in GFRP. Journal of IntelligentMaterial Systems and Structures 10(10):836–841.

Zhang, Y., A. A. Broekhuis, and F. Picchioni. 2009.Thermally self-healing polymeric materials: The next stepto recycling thermoset polymers? Macromolecules42(6):1906–1912.

Zhang, M.Q., and M. Z. Rong. 2011. Self-Healing Polymersand Polymer Composites: John Wiley & Sons, Hoboken, NJ.For cataloging purposes only

12 Chapter 12 - Preparation andProcessing of Magnesium Alloys

Al-Samman, T. 2009. Comparative study of the deformationbehavior of hexagonal magnesium-lithium alloys and aconventional magnesium AZ31 alloy. Acta Materialia 57:2229–42.

Al-Samman, T. and G. Gottstein. 2008. Room temperatureformability of a magnesium AZ31 alloy: Examining the roleof texture on the deformation mechanisms. Materials Scienceand Engineering A 488: 406–14.

Avedesian, M.M. and H. Baker. 1999. ASM SpecialtyHandbook-Magnesium and Magnesium Alloys. Materials Park,Ohio: ASM International.

Azeem, M.A., A. Tewari, S. Mishra, S. Gollapudi, and U.Ramamurty. 2010. Development of novel grain morphologyduring hot extrusion of magnesium AZ21 alloy. ActaMaterialia 58: 1495–502.

Azushima, A., R. Kopp, and A. Korhonen et al. 2008. Severeplastic deformation (SPD) processes for metals.Manufacturing Technology 57: 716–35.

Barnett, M.R. 2007. Twinning and the ductility of magnesiumalloys: Part II. “Contraction” twins. Materials Scienceand Engineering A 464: 8–16.

Barnett, M.R. 2007. Twinning and the ductility of magnesiumalloys: Part I: “Tension” twins. Materials Science andEngineering A 464: 1–7.

Barnett, M.R., A. Sullivan, N. Stanford, N. Ross, and A.Beer. 2010. Texture selection mechanisms in uniaxiallyextruded magnesium alloys. Scripta Materialia 63: 721–4.

Belyakov, A., W. Gao, H. Miura, and T. Sakai. 1998.Strain-induced grain evolution in polycrystalline copperduring warm deformation. Metallurgical and MaterialsTransactions A 29: 2957–65.

Bohlen, J., J. Swiostek, and W.H. Sillekens et al. 2005.Process and alloy development for hydrostatic extrusion ofmagnesium: The European Community research projectMAGNEXTRUSCO. In 2005 TMS Annual Meeting, February 13,2005–February 17, 2005. San Francisco, CA, United States:Minerals, Metals and Materials Society.

Cheng, Y., Z. Chen, W. Xia, and T. Zhou. 2008. Improvementof drawability at room temperature in AZ31 magnesium alloysheets processed by equal channel angular rolling. Journalof Materials Engineering and Performance 17: 15–9.

Chino, Y., M. Kado, and M. Mabuchi. 2008. Enhancement oftensile ductility and stretch formability of magnesium byaddition of 0.2 wt%(0.035 at%)Ce. Materials ScienceEngineering A 494: 343–9.

Chun, Y.B., J. Geng, and N. Stanford et al. 2011.Processing and properties of Mg-6Gd-1Zn-0.6Zr: Part 1—Recrystallisation and texture development. MaterialsScience and Engineering A 528: 3653–8.

del Valle, J.A., F. Carreno, and O.A. Ruano. 2006. In�uenceof texture and grain size on work hardening and ductilityin magnesium-based alloys processed by ECAP and rolling.Acta Materialia 54: 4247–59.

del Valle, J.A., M.T. Pérez-Prado, and O.A. Ruano. 2003.Texture evolution during large-strain hot rolling of theMg AZ61 alloy. Materials Science and Engineering A 355:68–78.

del Valle, J.A., M.T. Pérez-Prado, and O.A. Ruano. 2005.Accumulative roll bonding of a Mg-based AZ61 alloy.Materials Science and Engineering A 410–411: 353–7.

Emley, E.F. 1966. Principles of Magnesium Technology. NewYork: Pergamon Press.

Fatemi-Varzaneh, S.M., A. Zarei-Hanzaki, and M. Haghshenas.2009. The room temperature mechanical properties ofhot-rolled AZ31 magnesium alloy. Journal of Alloys andCompounds 475: 126–30.

Ferkel, H. and B.L. Mordike. 2001. Magnesium strengthenedby SiC nanoparticles. Materials Science and Engineering A298: 193–9.

Figueiredo, R.B. and T.G. Langdon. 2006. The development ofsuperplastic ductilities and microstructural homogeneityin a magnesium ZK60 alloy processed by ECAP. MaterialsScience and Engineering A 430: 151–6.

Flemings, M. 1991. Behavior of metal alloys in thesemisolid state. Metallurgical and Materials TransactionsA 22: 957–81.

Friedrich, H.E. and B.L. Mordike. 2006. MagnesiumTechnology-Metallurgy, Design Data, Applications. Berlin:Springer.

Furukawa, M., Y. Iwahashi, Z. Horita, M. Nemoto, and T.G.Langdon. 1998. The shearing characteristics associated withequal-channel angular pressing. Materials Science andEngineering A 257: 328–32.

Ghomashchi, M.R. and A. Vikhrov. 2000. Squeeze casting: Anoverview. Journal of Materials Processing Technology 101:1–9.

Guo, Q., H.G. Yan, Z.H. Chen, and H. Zhang. 2007. Grainre�nement in as-cast AZ80 Mg alloy under large straindeformation. Materials Characterization 58: 162–7.

Gupta, M. and N.M.L. Sharon. 2011. Magnesium, MagnesiumAlloys, and Mangnesium Composites. Hoboken, New Jersey:John Wiley & Sons, Inc. For cataloging purpos s nly

Han, L., H. Hu, D.O. Northwood, and N. Li. 2008.Microstructure and nano-scale mechanical behavior of Mg-Aland Mg-Al-Ca alloys. Materials Science and Engineering A473: 16–27.

Hassan, S.F. and M. Gupta. 2002. Development of a novelmagnesium-copper based composite with improved mechanicalproperties. Materials Research Bulletin 37: 377–89.

Hassan, S.F. and M. Gupta. 2003. Development of highstrength magnesium-copper based hybrid composites withenhanced tensile properties. Materials Science andTechnology 19: 253–9.

Hu, H. 1998. Squeeze casting of magnesium alloys and theircomposites. Journal of Materials Science 33: 1579–89.

Huang, X., K. Suzuki, A. Watazu, I. Shigematsu, and N.Saito. 2008. Microstructure and texture of Mg-Al-Zn alloyprocessed by differential speed rolling. Journal of Alloysand Compounds 457: 408–12.

Ibrahim, I.A., F.A. Mohamed, and E.J. Lavernia. 1991.Particulate reinforced metal matrix composites—A review.Journal of Materials Science 26: 1137–56.

Iwahashi, Y., J. Wang, Z. Horita, M. Nemoto, and T.G.Langdon. 1996. Principle of equal-channel angular pressingfor the processing of ultra-�ne grained materials. Scripta

Materialia 35: 143–6.

Jiang, L., J.J. Jonas, R.K. Mishra et al. 2007. Twinningand texture development in two Mg alloys subjected toloading along three different strain paths. Acta Materialia55: 3899–910.

Jiang, Q.C., H.Y. Wang, J.G. Wang, Q.F. Guan, and C.L. Xu.2003. Fabrication of TiCp/Mg composites by the thermalexplosion synthesis reaction in molten magnesium. MaterialsLetters 57: 2580–3.

Kim, W.J., C.W. An, Y.S. Kim, and S.I. Hong. 2002.Mechanical properties and microstructures of an AZ61 MgAlloy produced by equal channel angular pressing. ScriptaMaterialia 47: 39–44.

Kim, W.J., S.I. Hong, and Y.S. Kim et al. 2003. Texturedevelopment and its effect on mechanical properties of anAZ61 Mg alloy fabricated by equal channel angular pressing.Acta Materialia 51: 3293–307.

Kim, W.J., H.G. Jeong, and H.T. Jeong. 2009. Achieving highstrength and high ductility in magnesium alloys usingsevere plastic deformation combined with low-temperatureaging. Scripta Materialia 61: 1040–3.

Kim, W.J., J.B. Lee, W.Y. Kim, H.T. Jeong, and H.G. Jeong.2007. Microstructure and mechanical properties of Mg-Al-Znalloy sheets severely deformed by asymmetrical rolling.Scripta Materialia 56: 309–12.

Kim, S.-H., B.-S. You, C. Dong Yim, and Y.-M. Seo. 2005.Texture and microstructure changes in asymmetrically hotrolled AZ31 magnesium alloy sheets. Materials Letters 59:3876–80.

Kirkwood, D.H. 2008. Encyclopedia of Materials. New York:Elsevier.

Koh, H., T. Sakai, H. Utsunomiya, and S. Minamiguchi. 2007.Deformation and texture evolution during highspeed rollingof AZ31 magnesium sheets. Materials Transactions 48:2023–7.

Koike, J. 2005. Enhanced deformation mechanisms byanisotropic plasticity in polycrystalline Mg alloys atroom temperature. Metallurgical and Materials TransactionsA 36: 1689–96.

Lavernia, E.J. and Y. Wu. 1996. Spray Atomization andDeposition. New York: John Wiley & Sons, Ltd.

Lee, S.H., Y. Saito, N. Tsuji, H. Utsunomiya, and T. Sakai.2002. Role of shear strain in ultragrain re�nement byaccumulative roll-bonding (ARB) process. Scripta Materialia46: 281–5.

Li, H., E. Hsu, J. Szpunar, H. Utsunomiya, and T. Sakai.2008. Deformation mechanism and texture and microstructureevolution during high-speed rolling of AZ31B Mg sheets.Journal of Materials Science 43: 7148–56.

Liu, J. and Z. Cui. 2009. Hot forging process design andparameters determination of magnesium alloy AZ31B spurbevel gear. Journal of Materials Processing Technology 209:5871–80.

Liu, S.F., L.Y. Liu, and L.G. Kang. 2008. Re�nement role ofelectromagnetic stirring and strontium in AZ91 magnesiumalloy. Journal of Alloys and Compounds 450: 546–50.

Liu, S.F., B. Li, X.H. Wang, W. Su, and H. Han. 2009.Re�nement effect of cerium, calcium and strontium in AZ91magnesium alloy. Journal of Materials Processing Technology209: 3999–4004.

Lloyd, D.J. 1994. Particle reinforced aluminium andmagnesium matrix composites. International MaterialsReviews 39: 1–23.

Mabuchi, M., H. Iwasaki, K. Yanase, and K. Higashi. 1997.Low temperature superplasticity in an AZ91 magnesium alloyprocessed by ECAE. Scripta Materialia 36: 681–6.

Mabuchi, M., K. Ameyama, H. Iwasaki, and K. Higashi. 1999.Low temperature superplasticity of AZ91 magnesium alloywith non-equilibrium grain boundaries. Acta Materialia 47:2047–57.

Martin, É., L. Capolungo, L. Jiang, and J.J. Jonas. 2010.Variant selection during secondary twinning in Mg-3%Al.Acta Materialia 58: 3970–83.

Minkoff, I. 1986. Solidi�cation and Cast Structure. NewYork: Wiley.

Mishra, R.K., A.K. Gupta, P.R. Rao et al. 2008. In�uence ofcerium on the texture and ductility of magnesiumextrusions. Scripta Materialia 59: 562–5.

Nayeb-Hashemi, A.A. and J.B. Clark. 1998. Phase Diagrams ofBinary Magnesium Alloys. Metals Park, OH: ASMInternational.

Pérez-Prado, M.T., J.A. del Valle, and O.A. Ruano. 2004.Grain re�nement of Mg-Al-Zn alloys via accumulative rollbonding. Scripta Materialia 51: 1093–7.

Pérez-Prado, M.T., J.A. del Valle, and O.A. Ruano. 2005.Achieving high strength in commercial Mg cast alloysthrough large strain rolling. Materials Letters 59:3299–303.

Pérez-Prado, M.T., J.A. del Valle, J.M. Contreras, and O.A.Ruano. 2004. Microstructural evolution during large strainhot rolling of an AM60 Mg alloy. Scripta Materialia 50:661–5.

Ping, W., J. Li, and M. Qun. 2008. Effects of gadolinium onthe microstructure and corrosion resistance properties ofZK60 magnesium alloy. Rare Metal Materials and Engineering37: 1056–9.

Richert, M., H.J. McQueen, and J. Richert. 1998. Microbandformation in cyclic extrusion compression of aluminum.Canadian Metallurgical Quarterly 37: 449–57.

Richert, M., Q. Liu, and N. Hansen. 1999. Microstructuralevolution over a large strain range in aluminium deformedby cyclic-extrusion-compression. Materials Science andEngineering A 260: 275–83.

Saito, Y., N. Tsuji, H. Utsunomiya, T. Sakai, and R.G.Hong. 1998. Ultra-�ne grained bulk aluminum produced byaccumulative roll-bonding (ARB) process. Scripta Materialia39: 1221–7.

Saito, Y., H. Utsunomiya, N. Tsuji, and T. Sakai. 1999.Novel ultra-high straining process for bulk materials—Development of the accumulative roll-bonding (ARB) process.Acta Materialia 47: 579–83.

Segal, V.M. 2002. Severe plastic deformation: Simple shearversus pure shear. Materials Science and Engineering A338: 331–44.

Sillekens, W.H., J.A.F.M. Schade van Westrum, A.J. DenBakker, and P.J. Vet. 2003. Hydrostatic extrusion ofmagnesium: Process mechanics and performance. In Thermec

2003 Processing and Manufacturing of Advanced Materials,July 7, 2003–July 11, 2003. Madrid, Spain: Trans TechPublications Ltd.

Sitdikov, O., T. Sakai, A. Goloborodko, and H. Miura. 2004.Grain fragmentation in a coarse-grained 7475 Al alloyduring hot deformation. Scripta Materialia 51: 175–9.

Spencer, D., R. Mehrabian, and M. Flemings. 1972.Rheological behavior of Sn-15 pct Pb in the crystallizationrange. Metallurgical and Materials Transactions B 3:1925–32.

Stanford, N. and M.R. Barnett. 2009. Effect of particles onthe formation of deformation twins in a magnesiumbasedalloy. Materials Science and Engineering: A 516: 226–34.

Suzuki, M., H. Sato, K. Maruyama, and H. Oikawa. 2001.Creep deformation behavior and dislocation substructures ofMg-Y binary alloys. Materials Science and Engineering A319–321: 751–5.

Swiostek, J., J. Göken, D. Letzig, and K. Kainer. 2006.Forming of magnesium alloys at 100°C by hydrostaticextrusion. Journal of Materials Engineering and Performance15: 705–11.

Takuda, H., T. Enami, K. Kubota, and N. Hatta. 2000. Theformability of a thin sheet of Mg-8.5Li-1Zn alloy. Journalof Materials Processing Technology 101: 281–6.

Thakur, S.K. and B.K. Dhindaw. 2001. The in�uence ofinterfacial characteristics between SiCp and Mg/Al metalmatrix on wear, coef�cient of friction and microhardness.Wear 247: 191–201.

Tissier, A., D. Apelian, and G. Regazzoni. 1990. Magnesiumrheocasting: A study of processing-microstructureinteractions. Journal of Materials Science 25: 1184–96.

TMS, U.S. 2011a. Magnesium Extrusion Alloys.http://iweb.tms.org/forum/categories.aspx?catid=364&entercat=y.

TMS, U.S. 2011b. Magnesium Forging Alloys

TMS, U.S. 2011c. Magnesium Sheet and Plate Alloyshttp://iweb.tms.org/forum/categories.aspx?catid=364&entercat=y.

Tsuji, N., Y. Saito, S.-H. Lee, and Y. Minamino. 2003. ARB(accumulative roll-bonding) and other new techniques toproduce bulk ultra�ne grained materials. AdvancedEngineering Materials 5: 338–44.

Vaidya, A.R. and J.J. Lewandowski. 1996. Effects of SiCpsize and volume fraction on the high cycle fatiguebehavior of AZ91D magnesium alloy composites. MaterialsScience and Engineering A 220: 85–92.

Wang, H.Y., Q.C. Jiang, Y. Wang, B.X. Ma, and F. Zhao.2004. Fabrication of TiB2 particulate reinforced magnesiummatrix composites by powder metallurgy. Materials Letters58: 3509–13.

Wang, M., R. Xin, B. Wang, and Q. Liu. 2011. Effect ofinitial texture on dynamic recrystallization of AZ31 Mgalloy during hot rolling. Materials Science and EngineeringA 528: 2941–51.

Wang, Q.D., Y.J. Chen, J.B. Lin, L.J. Zhang, and C.Q. Zhai.2007. Microstructure and properties of magnesium alloyprocessed by a new severe plastic deformation method.Materials Letters 61: 4599–602.

Watanabe, H., T. Mukai, and K. Ishikawa. 2004. Differentialspeed rolling of an AZ31 magnesium alloy and the resultingmechanical properties. Journal of Materials Science 39:1477–80.

Watanabe, H., T. Mukai, and K. Ishikawa. 2007. Effect oftemperature of differential speed rolling on roomtemperature mechanical properties and texture in an AZ31magnesium alloy. Journal of Materials ProcessingTechnology 182: 644–7. For c t logi g p rp ses nly

Watanabe, H., T. Mukai, and K. Ishikawa et al. 2001.Superplasticity of a Particle-Strengthened WE43 MagnesiumAlloy. Materials Transactions 42: 157–62.

Watari, H., K. Davey, M.T. Rasgado, T. Haga, and S. Izawa.2004. Semi-solid manufacturing process of magnesium alloysby twin-roll casting. Journal of Materials ProcessingTechnology 155–156: 1662–7.

Wong, W.L.E. and M. Gupta. 2007. Development of Mg/Cunanocomposites using microwave assisted rapid sintering.Composites Science and Technology 67: 1541–52.

Xie, S., Y. He, G. Huang, M. Geng, and Y. Zhang. 2008.

Study on semi-solid continuous roll-casting strips ofAZ91D magnesium alloy. In Semi-Solid Processing of Alloysand Composites 10—Selected, peer reviewed papers from the10th International Conference on Semi-Solid Processing ofAlloy and Composites, S2P 2008, September 16,2008–September 18, 2008. Aachen, Germany: Trans TechPublications Ltd.

Xie, S., H. Yang, L. Li, and G. Huang 2006. Numericalsimulation of semi-solid magnesium alloy in continuousroll-casting process. In 9th International Conference onSemi-Solid Processing of Alloys and Composites, S2P 2006,September 11, 2006–September 13, 2006. Busan, Korea,Republic of: Trans Tech Publications Ltd.

Xin, Y., M. Wang, Z. Zeng, G. Huang, and Q. Liu. 2011.Tailoring the texture of magnesium alloy by twinningdeformation to improve the rolling capability. ScriptaMaterialia 64: 986–9.

Xing, J., X. Yang, H. Miura, and T. Sakai. 2008. Mechanicalproperties of magnesium alloy AZ31 after severe plasticdeformation. Materials Transactions 49: 69–75.

Yan, J., Y. Sun, F. Xue, S. Xue, and W. Tao. 2008.Microstructure and mechanical properties in castmagnesiumneodymium binary alloys. Materials Science andEngineering A 476: 366–71.

Yang, P., Y. Yu, L. Chen, and W. Mao. 2004. Experimentaldetermination and theoretical prediction of twinorientations in magnesium alloy AZ31. Scripta Materialia50: 1163–8.

Yim, C., B. You, R. Jang, and S. Lim. 2006. Effects of melttemperature and mold preheating temperature on the �uidityof Ca containing AZ31 alloys. Journal of Materials Science41: 2347–50.

Yoo, M. 1981. Slip, twinning, and fracture in hexagonalclose-packed metals. Metallurgical and MaterialsTransactions A 12: 409–18.

Yu, Y. and E. Wang. 2009. The effects of cold extrusion ongrain size re�nement and plasticity for magnesium alloy.International Journal of Modern Physics B 23: 821–5.

Zarandi, F., G. Seale, R. Verma, E. Essadiqi, and S. Yue.2008. Effect of Al and Mn additions on rolling anddeformation behavior of AZ series magnesium alloys.

Materials Science and Engineering: A 496: 159–68.

Zhang, Y., M. Geng, L. Cheng, and Y. Wang. 2008. In�uenceof processing parameters on microstructure of castingrolling semi-solid AZ91D magnesium alloy. In Semi-SolidProcessing of Alloys and Composites 10—Selected, peerreviewed papers from the 10th International Conference onSemi-Solid Processing of Alloy and Composites, S2P 2008,September 16, 2008–September 18, 2008. Aachen, Germany:Trans Tech Publications Ltd.

Zhao, Y., F. Pan, J. Peng, W. Wang, and S. Luo. 2010.Effect of neodymium on the as-extruded ZK20 magnesiumalloy. Journal of Rare Earths 28: 631–5.

Zherebtsov, S.V., G.A. Salishchev, R.M. Galeyev et al.2004. Production of submicrocrystalline structure inlarge-scale Ti-6Al-4V billet by warm severe deformationprocessing. Scripta Materialia 51: 1147–51.

13 Chapter 13 - Fatigue of MagnesiumAlloys

1. Gillett, N. P., Arora, V. K., Zickfeld, K., Marshall, S.J., and W. J. Merry�eld. 2011. Ongoing climate changefollowing a complete cessation of carbon dioxide emissions.Nat. Geosci. 4:83–87.

2. Min, S. K., Zhang, X. B., Zwiers, F. W., and G. C.Hegerl. 2011. Human contribution to more-intenseprecipitation extremes. Nature 470:378–81.

3. Schiermeier, Q. 2011. Increased �ood risk linked toglobal warming. Nature 470:316.

4. Pall, P., Aina, T., Stone, D. A., Scott, P. A., Nozawa,T., Hilberts, A. G. J., Lohmann, D., and M. R. Allen.2007. Anthropogenic greenhouse gas contribution to �oodrisk in England and Wales in autumn 2000. Nature470(7334):382–85.

5. Hassan, S. F., and M. Gupta. 2004. Development of highperformance magnesium nanocomposites using solidi�cationprocessing route. Mat. Sci. Tech. 20:1383–88.

6. Pollock, T. M. 2010. Weight loss with magnesium alloys.Science 328:986–87.

7. Wise, M., Calvin, K., Thomson, A., Clarke, L., Lamberty,B. B., Sands, R., Smith, S. J., Janetos, A., and J.Edmonds. 2009. Implications of limiting CO 2concentrations for land use and energy. Science324:1183–86.

8. Kump, L. R. 2002. Reducing uncertainty about carbondioxide as a climate driver. Nature 419:188–90.

9. Agnew, W. G. 1974. Reducing automotive emissions.Science 183(4122):254–56.

10. Mordike, B. L., and T. Ebert. 2001. Magnesium:Properties-applications-potential. Mater. Sci. Eng. A.302(1):37–45.

11. Eliezer, D., Aghion, E., and F. H. Froes. 1998.Magnesium science, technology and applications. Adv.Perfor. Mater. 5:201–12.

12. Kainer, K. U. 2003. Magnesium—Alloys and Technology.Wiley-VCH, Cambridge.

13. Luo, A. A. 2004. Recent magnesium alloy development forelevated temperature applications. Int. Mater. Rev.49(1):13–30.

14. Ke, W., Han, E. H., Han, Y. F., Kainer K., and A. A.Luo. 2005. Mechanical properties and creep behavior ofMg-Al-Ca alloys. Proceedings of International Conference onMagnesium—Science, Technology and Applications. Beijing,China. Mater. Sci. Forum. 488–89.

15. United States Automotive Materials Partnership (USAMP).2006. Magnesium vision 2020: A North American automotivestrategic vision for magnesium.

16. Duffy, L. 1996. Magnesium alloys—An introduction.Mater. World 4(3):127–30.

17. Pantelakis, Sp. G., Alexopoulos, N. D., and A. N.Chamos. 2007. Mechanical performance evaluation of castmagnesium alloys for automotive and aeronauticalapplications. J. Eng. Mater. Tech. 129:422–30.

18. Göken, J., Bohlen, J., Hort, N., Letzig, D., andKainer, K.U. 2003. New development in magnesium technologyfor lightweight structures in transportation industries.Mater. Sci. Forum 426–432(1):153–60.

19. Zeng, R. C., Han, E. H., and W. Ke. 2005. Fatigue andcorrosion fatigue of magnesium alloys. Mater. Sci. Forum488–489:721–24.

20. Makar, G.L., and Kruger, J. 1993. Corrosion ofmagnesium. Int. Mater. Rev. 38(3):138–53.

21. Ghali, E., Dietzel, W., and K. U. Kainer. Testing ofgeneral and localized corrosion of magnesium alloys: Acritical review. J. Mater. Eng. Per. 13(5):517–29.

22. Pekguleryuz, M. O., and A. A. Kaya. 2004. Creepresistant magnesium alloys for powertrain applications. InMagnesium: Proceedings of the 6th International ConferenceMagnesium Alloys and Their Applications. Edited by K. U.Kainer. Wiley-Vch Verlag GmbH, Weinheim, Germany. pp.74–93.

23. Nicholas, J. 2005. High performance magnesium. Adv.Mater. Process. 163(9):65–67.

24. ASM International. The Materials Information Society.

2011. Research improves the bolted joints in airplanes.

25. Wang, Q., Zhang, Z. M., Zhang, X., and G. J., Li. 2010.New extrusion process of Mg alloy automobile wheels.Trans. Nonfer. Metal. Soc. China 20(2):599–03.

26. Potzies, C., and K. U. Kainer. 2004. Fatigue ofmagnesium alloys. Adv. Eng. Mater. 6(5):281–89.

27. Begum, S., Chen, D. L., Xu, S., and A. A. Luo. 2009.Effect of strain ratio and strain rate on low cyclefatigue behavior of AZ31 wrought magnesium alloy. Mater.Sci. Eng. A. 517:334–43.

28. Begum, S., Chen, D. L., Xu, S., and A. A. Luo. 2009.Low cycle fatigue properties of an extruded AZ31 magnesiumalloy. Int. J. Fatigue 31:726–35.

29. Lin, X. Z., and D. L. Chen. 2008. Strain controlledcyclic deformation behavior of an extruded magnesium alloy.Mater. Sci. Eng. A. 496:106–13.

30. Brown, D. W., Jain, A., Agnew, S., and R. B. Clausen.2007. Twinning and detwinning during cyclic deformation ofMg alloy AZ31B. Mater. Sci. Forum 539–543:3407–13.

31. Noster, U., and B. Scholtes. 2003. Isothermalstrain-controlled quasi-static and cyclic deformationbehavior of magnesium wrought alloy AZ31. Mater. Res. Adv.Tech. (Zeitschrift fuer Metallkunde). 94(5):559–63.

32. Hasegawa, S., Tsuchida, Y., Yano, H., and M. Matsui.2007. Evaluation of low cycle fatigue life in AZ31magnesium alloy. Int. J. Fatigue 29:1839–45.

33. Begum, S., Chen, D. L., Xu, S., and A. A. Luo. 2008.Strain-controlled low-cycle fatigue properties of a newlydeveloped extruded magnesium alloy. Metall. Mat. Trans. A.39:3014–26.

34. Fan, C. L., Chen, D. L., and A. A. Luo. 2009.Dependence of the distribution of deformation twins onstrain amplitudes in an extruded magnesium alloy aftercyclic deformation. Mater. Sci. Eng. A. 519:38–45.

35. Wu, L., Jain, A., Brown, D. W., Stoica, G. M, Agnew, S.R., Clausen, B., Fielden, D. E., and P. K. Liaw. 2008.Twinning-detwinning behavior during the strain-controlledlow-cycle fatigue testing of a wrought magnesium alloy,ZK60A. Acta Mater. 56:688–95.

36. Wu, L., Agnew, S. R., Brown, D.W., Stoica, G. M,Clausen, B., Jain, B., Fielden, D. E., and P. K. Liaw.2008. Internal stress relaxation and load redistributionduring the twinning-detwinning-dominated cyclicdeformation of a wrought magnesium alloy, ZK60A. ActaMater. 56:3699–07.

37. Eisenmeier, G., Holzwarth, B., Hoppel, H. W., and H.Mughrabi. 2001. Cyclic deformation and fatigue behaviourof the magnesium alloy AZ91. Mater. Sci. Eng. A.319–321:578–82.

38. Chen, L., Shen, J., Wu, W. Li, F., Wang, Y., and Z.Liu. 2005. Low-cycle fatigue behavior of magnesium alloyAZ91. Mater. Sci. Forum 488–489:725–28. For catalogingpurp s o ly

39. Horstemeyer, M. F., Yang, N., Gall, K., McDowell, D.L., Fan, J., and P. M. Gullett. 2004. High cycle fatigueof a die cast AZ91E-T4 magnesium alloy. Acta Mater.52:1327–36.

40. Liu, Z., Xu, Y. Y., Wang, Z.G., Wang, Y., and Z.Y. Liu.2000. Low cycle fatigue behavior of AZ91HP alloy in ashigh pressure die-casting. Acta Metall. Sin. (EnglishLetters) 13:961–66.

41. Islamgaliev, R. K., Kulyasova, O. B., Mingler, B.,Zehetbauer, M., and A. Minkow. 2008. Structure and fatigueproperties of the Mg alloy AM60 processed by ECAP. Mater.Sci. Forum 584–586:803–08.

42. Kulyasova, O., Islamgaliev, R., Mingler, B., and M.Zehetbauer. 2009. Microstructure and fatigue properties ofthe ultra�ne-grained AM60 magnesium alloy processed byequal-channel angular pressing. Mater. Sci. Eng. A.503:176–80.

43. El Kadiri, H., Horstemeyer, M. F., Jordon, J. B., andY. Xue. 2008. Fatigue crack growth mechanisms inhigh-pressure die-cast magnesium alloys. Metall. Mat.Trans. A. 39:190–05.

44. Horstemeyer, M. F., Yang, N. K. Gall, McDowell, D.,Fan, J., and P. Gullett. 2002. High cycle fatiguemechanisms in a cast AM60B magnesium alloy. Fatigue Fract.Eng. Mater. Struct. 25:1045–56.

45. Liu, Z., Ji, H., Lin, L., Chen, L., Wu, W., and L.

Yang. 2007. Cyclic deformation behavior and potentialautomobile application of magnesium die casting alloys AZ91and AM50, Mater. Sci. Forum 539–543:1626–31.

46. El Kadiri, H., Xue, Y., Horstemeyer, M. F., Jordon, J.B., and P.T. Wang. 2006. Identi�cation and modeling offatigue crack growth mechanisms in a die-cast AM50magnesium alloy. Acta Mater. 54(19):5061–76.

47. Liu, Z., Wang, Z. G., Wang, Y., and Z. Y. Liu. 1999.Cyclic deformation behavior of high pressure die castingalloy AM50. J. Mater. Sci. Lett. 18:1567–69.

48. Unigovski, Y., Eliezer, A., Abramov, E., Snir, Y., andE. M. Gutman. 2003. Corrosion fatigue of extrudedmagnesium alloys. Mater. Sci. Eng. A. 360:132–39.

49. Barbagallo, S., Laukli, H. I., Lohne, O., and E. Cerri.2004. Divorced eutectic in a HPDC magnesiumaluminum alloy.J. Alloys Comp. 378:226–32.

50. Moreno, I. P., Nandy, T. K., Jones, J. W., Allison, J.E., and T. M. Pollock. 2001. Microstructuralcharacterization of a die-cast magnesium-rare earth alloy.Scripta Mater. 45(12):1423–29.

51. Cai, J., Ma, G. C., Liu, Z., Zhang, H. F., and Z. Q.Hu. 2006. In�uence of rapid solidi�cation on themicrostructure of AZ91HP alloy. J. Alloys Comp.422(1–2):92–96.

52. Wang, R. M., Eliezer, A., and E. M. Gutman. 2003. Aninvestigation on the microstructure of an AM50 magnesiumalloy. Mater. Sci. Eng. A. 355:201–07.

53. ASM Handbook. 2004. Metallography and Microstructures.ASM, Materials Park, OH. Vol. 9, p. 813, Figure 27.

54. He, S. M., Zeng, X. Q., Peng, L. M., Gao, X., Nie, J.F., and W. J., Ding. 2007. Microstructure andstrengthening mechanism of high strength Mg-10Gd-2Y-0.5Zralloy. J. Alloys Comp. 427(1–2):316–23.

55. Lü, Y. Z., Wang, Q. D., Zeng, X. Q., Ding, W. J., Zhai,C. Q., and Y. P., Zhu. 2000. Effects of rare earths on themicrostructure, properties and fracture behavior of Mg-Alalloys. Mater. Sci. Eng. A. 278(1–2):66–76.

56. He, S. M., Peng, L. M., Zeng, X. Q., Ding, W. J., andY. P., Zhu. 2006. Comparison of the microstructure and

mechanical properties of a ZK60 alloy with and without 1.3wt.% gadolinium addition. Mater. Sci. Eng. A.433(1–2):175–81.

57. Li, Q., Wang, Q. D., Wang, Y. X., Zeng, X. Q., and W.J., Ding. 2007. Effect of Nd and Y addition onmicrostructure and mechanical properties of as-castMg-Zn-Zr alloy. J. Alloys Comp. 427(1–2):115–23.

58. Patel, H. A., Chen, D. L., Bhole, S. D., and K.Sadayappan. 2010. Microstructure and tensile properties ofthixomolded magnesium alloys. J. Alloys Comp.496(1-2):140–48.

59. Zhang, Y. F., Liu, Y. B., Zhang, Q. Q., Cao, Z. Y.,Cui, X. P., and Y. Wang. 2007. Microstructural evolution ofthixomolded AZ91D magnesium alloy with process parametersvariation. Mater. Sci. Eng. A. 444:251–56.

60. Fan, Z. 2002. Semisolid metal processing. Int. Mater.Rev. 47:49–86.

61. Cao, X., Jahazi, M., Immarigeon, J. P., and W. Wallace.2006. A review of laser welding techniques for magnesiumalloys. J. Mater. Proc. Tech. 171(2):188–04.

62. Fairman, M., Afrin, N., Chen, D. L., Cao, X. J., and M.Jahazi. 2007. Microstructural evaluation of friction stirprocessed AZ31B-H24 magnesium alloy. Can. Metall. Quart.46:425–32.

63. Afrin, N., Chen, D. L., Cao, X., and M. Jahazi. 2007.Strain hardening behaviour of a friction stir weldedmagnesium alloy. Scripta Mater. 57:1004–07.

64. Afrin, N., Chen, D. L., Cao, X., and M. Jahazi. 2008.Microstructure and tensile properties of friction stirwelded AZ31B magnesium alloy. Mater. Sci. Eng. A.472:179–86.

65. Chowdhury, S. M., Chen, D. L., Bhole, S. D., and X.Cao. 2010. Tensile properties of a friction stir weldedmagnesium alloy: Effect of pin tool thread orientation andweld pitch. Mater. Sci. Eng. A. 527(21–22):6064–75.

66. Chowdhury, S. M., Chen, D. L., Bhole, S.D., Powidajko,E., Weckman, D. C., and Y. Zhou. 2011. Microstructure andmechanical properties of �ber-laser-welded anddiode-laser-welded AZ31 magnesium alloy. Metall. Mater.Trans. A. 42A:1974–89.

67. Chowdhury, S. H., Chen, D. L., S. D. Bhole, S.D.,Powidajko, E., Weckman, D. C., and Y. Zhou. 2012. Fiberlaser welded AZ31 magnesium alloy: Effect of welding speedon microstructure and mechanical properties. Metall.Mater. Trans. A. 43(6):2133–47.

68. Mohanty, P. S., and J. Mazumder. 1998. Solidi�cationbehavior and microstructural evolution during laserbeam-material interaction. Metall. Mater. Trans. B.29B:1269–79.

69. Kou, S., and Y. Le. 1985. Alternating grain orientationand weld solidi�cation cracking. Metall. Trans. A.16(10):1887–96.

70. Rappaz, M., Vitek, J. M., David, S. A., and L.A.Boatner. 1993. Microstructural formation in longitudinalbicrystal welds. Metall. Trans. A. 24:1433–46.

71. Somboonsuk, K., Mason, J. T., and R. Trivedi. 1984.Interdendritic spacing: Part I. experimental studies.Metall. Trans. A. 15(6):967–75.

72. Wu, S. H., Huang, J. C., and Y. N. Wang. 2004.Evolution of microstructure and texture in Mg-Al-Zn alloysduring electron-beam and gas tungsten arc welding. Metall.Trans. A. 35(8):2455–69.

73. Liu, Y., Koch, J., Mazumder, J., and K. Shibata. 1994.Processing, microstructure, and properties of laserclad Nialloy FP-5 on Al alloy AA333. Metalls Trans. B.25(3):425–34.

74. Yajie, Q., Chen, Z., Yu, Z., Gong, X., and M. Li. 2008.Effects of heat input on microstructure and tensileproperties of laser welded magnesium alloy AZ31. Mater.Char. 59(10):1799–04.

75. Yu, L., Nakata, K., Yamamoto, N., and J. Liao. 2009.Texture and its effect on mechanical properties in �berlaser weld of a �ne-grained Mg alloy. Mater. Lett.63(11):870–72.

76. Yu, Z. H., Yan, H. G., Gong, X. S., Quan, Y. J., Chen,J. H., and Q. Chen. 2009. Microstructure and mechanicalproperties of laser welded wrought ZK21 magnesium alloy.Mat. Sci. Eng. A. 523(1–2):220–25.

77. Weisheit, A., Galun, R., and B. L. Mordike. 1998. CO 2

laser beam welding of magnesium-based alloys. Weld. J.77(4):148–54.

78. Zhao, H., and T. Debroy. 2001. Pore formation duringlaser beam welding of die-cast magnesium alloyAM60B—Mechanism and remedy. Weld. J. 80(8):204–10.

79. Sun, Z., Pan, D., and J. Wei. 2002. Comparativeevaluation of tungsten inert gas and laser welding of AZ31magnesium alloy. Sci. Tech. Weld. J. 7(6):343–51.

80. Padmanaban, G., and V. Balasubramanian. 2010. Fatigueperformance of pulsed current gas tungsten arc, frictionstir and laser beam welded AZ31B magnesium alloy joints.Mater. Des. 31(8):3724–32.

81. Sun, H., Song, G., and L. F. Zhang. 2008. Effects ofoxide activating �ux on laser welding of magnesium alloy.Sci. Tech. Weld. J. 13(4):305–11.

82. Liu, L. M. 2010. Welding and Joining of MagnesiumAlloys. Woodhead Publishing in Materials, Cambridge,England.

83. Ren, D. X., and L. M., Liu. 2010. Effect of adhesiveinduced gas on penetration and plasma behaviour in laserweld bonding of magnesium alloy. Mater. Res. Inno.14(5):405–09.

84. Liu, L. M., and X. F., Hao. 2010. Low-power laser/TIGhybrid welding process of magnesium alloy with �ller wire,Mater. Manuf. Process. 25(11):1213–18.

85. Song, G., Wang, P., and L. M. Liu. 2010. Study onac-PMIG welding of AZ31B magnesium alloy. Sci. Technol.Welding Joining 15(3):219–25.

86. Liu, L. M., Song, G., and M. L. Zhu. 2008. Low-powerlaser/arc hybrid welding behavior in AZ-based Mg alloys.Metall. Mater. Trans. A. 39(7):1702–11.

87. Xie, G. M., Ma, Z. Y., and L. Geng. 2008. Effect ofmicrostructural evolution on mechanical properties offriction stir welded ZK60 alloy. Mater. Sci. Eng. A.486(1-2):49–55.

88. Xie, G. M., Ma, Z. Y., Geng, L., and R.S. Chen. 2007.Microstructural evolution and mechanical properties offriction stir welded Mg-Zn-Y-Zr alloy. Mater. Sci. Eng. A.471(1–2):63–68.

89. Xiao, L., Liu, L., Chen, D. L., Esmaeili, S., and Y.Zhou. 2011. Resistance spot weld fatigue behavior anddislocation substructures in two different heats of AZ31magnesium alloy. Mater. Sci. Eng. A. 529:81–87.

90. Xiao, L., Liu, L., Zhou, Y., and S., Esmaeili. 2010.Resistance-spot-welded AZ31 magnesium alloys: Part I.Dependence of fusion zone microstructures on second-phaseparticles. Metall. Mater. Trans. A: Phy. Metall. Mater.Sci. 41(6):1511–22.

91. Liu, L., Xiao, L., Feng, J. C., Tian, Y. H., Zhou, S.Q., and Zhou, Y. 2010. Resistance spot welded AZ31magnesium alloys, part II: Effects of welding current onmicrostructure and mechanical properties. Metall. Mater.Trans. A: Phy. Metall. Mater. Sci. 41(10):2642–50.

92. Liu, L., Zhou, S. Q., Tian, Y. H., Feng, J. C., Jung,J. P., and Y. N., Zhou. 2009. Effects of surface conditionson resistance spot welding of Mg alloy AZ31. Sci. Tech.Welding Joining 14(4):356–61. For cat loging purp ses only

93. Patel, V. K., Bhole, S. D., and D. L. Chen. 2011.In�uence of ultrasonic spot welding on microstructure in amagnesium alloy. Scripta Mater. 65(10):911–14.

94. Kim, S. H., You, B. S., You, B. S., and C. D. Yim.2005. The effect of rolling conditions on themicrostructure and texture evolution of AZ31 Mg alloysheets. Mater. Forum 29:530–35.

95. Johnson, R. 2003. Friction stir welding of magnesiumalloys. Mater. Sci. Forum 419–422:365–70.

96. Wang, X. H., and K. Wang. 2006. Microstructure andproperties of friction stir butt-welded AZ31 magnesiumalloy. Mater. Sci. Eng. A. 431:114–17.

97. Park, S. H. C., Sato, Y. S., and H. Kokawa. 2003.Effect of micro-texture on fracture location in frictionstir weld of Mg alloy AZ61 during tensile test. ScriptaMater. 49:161–66.

98. Woo, W., Choo, H., Prime, M. B., Feng, Z., and B.Clausen. 2008. Microstructure, texture and residual stressin a friction-stir-processed AZ31B magnesium alloy. ActaMater. 56:1701–11.

99. Commin, L., Dumont, M., Masse, J. E., and L.

Barrallier. 2009. Friction stir welding of AZ31 magnesiumalloy rolled sheets: In�uence of processing parameters.Acta Mater. 57:326–34.

100. Seidel, T.U., and A. P. Reynolds. 2001. Visualizationof the material �ow in AA2195 friction-stir welds using amarker strain technique. Metall. Mater. Trans. A.32:2879–84.

101. Fratini, L., Buffa, G., Palmeri, D., Hua, J., and R.Shivpur. 2006. Material �ow in FSW of AA7075-T6 buttjoints: Continuous dynamic recrystallization phenomena. J.Eng. Mater. Tech. 128:228–35.

102. Xu, W. F., Liu, J. H., and D.L. Chen. 2011. Material�ow and core/multi-shell structures in a friction stirwelded aluminum alloy. J. Alloys Comp. 509(33):8449–54.

103. Cui, G. R., Ma, Z. Y., and S.X. Li. 2008. Periodicalplastic �ow pattern in friction stir processed Al-Mgalloy. Scripta Mater. 58:1082–85.

104. Schmidt, H. N. B., Dickerson, T. L., and J. H. Hattel.2006. Material �ow in butt friction stir welds inAA2024-T3. Acta Mater. 54:1199–09.

105. Mishra, R. S., and Z. Y. Ma. 2005. Friction stirwelding and processing. Mater. Sci. Eng. R. 50(1-2):1–78.

106. Çam, G. 2011. Friction stir welded structuralmaterials: Beyond Al-alloys. Int. Mater. Rev. 56(1):1–48.

107. Threadgilll, P. L., Leonard, A. J., Shercliff, H. R.,and P.J. Withers. 2009. Friction stir welding of aluminiumalloys. Int. Mater. Rev. 54(2):49–93.

108. Ma, Z. Y. 2008. Friction stir processing technology: Areview. Metall. Mater. Trans. A. 39 A(3):642–58.

109. Nandan, R., DebRoy, T., and H. K. D. H. Bhadeshia,2008. Recent advances in friction-stir weldingProcess,weldment structure and properties. Prog. Mater. Sci.53(6):980–23.

110. Chowdhury, S. M., Chen, D. L., Bhole, S. D., Cao, X.,Powidajko, E., Weckman, D. C., and Y. Zhou. 2010. Tensileproperties and strain hardening behavior of double-sidedarc welded and friction stir welded AZ31B magnesium alloy.Mater. Sci. Eng. A. 527(12):2951–61.

111. Cao, X., and M. Jahazi. 2009. Effect of welding speedon the quality of friction stir welded butt joints of amagnesium alloy. Mater. Des. 30:2033–42.

112. Esparza, J. A., Davis, W. C., Trillo, E. A., and L. E.Murr. 2002. Friction stir welding of magnesium alloy AZ31B. J. Mater. Sci. Lett. 21:917–20.

113. Jiang, L., Jonas, J. J., Luo, A. A., Sachdev, A. K.,and S. Godet. 2006. Twinning-induced softening inpolycrystalline AM30 Mg alloy at moderate temperatures.Scripta Mater. 54:771–75.

114. Barnett, M. R. 2007. Twinning and the ductility ofmagnesium alloys. Part I: “Tension” twins. Mater. Sci.Eng. A. 464(1–2):1–7.

115. Barnett, M. R. 2007. Twinning and the ductility ofmagnesium alloys. Part II. “Contraction” twins. Mater.Sci. Eng. A. 464(1–2):8–16.

116. Jain, A., and S. R., Agnew. 2007. Modeling thetemperature dependent effect of twinning on the behaviorof magnesium alloy AZ31B sheet. Mater. Sci. Eng. A.462(1–2):29–36.

117. Bohlen, J., Nürnberg, M. R., Senn, J. W., Letzig, D.,and S. R. Agnew. 2007. The texture and anisotropy ofmagnesium-zinc-rare earth alloy sheets. Acta Mater.55(6):2101–12.

118. Wang, Y. N., and J. C., Huang. 2007. The role oftwinning and untwinning in yielding behavior in hotextrudedMg-Al-Zn alloy. Acta Mater. 55(3):897–05.

119. Zeng, R., Han, E., Ke, W., Dietzel, W., Kainer, K. U.,and A. Atrens. 2010. In�uence of microstructure on tensileproperties and fatigue crack growth in extruded magnesiumalloy AM60. Int. J. Fatigue 32:411–19.

120. Yan, C., Ma, W., Burg, V., and M. W. Chen. 2007.Experimental and numerical investigation on ductilebrittlefracture transition in a magnesium alloy. J. Mater. Sci.42:7702–07.

121. Albinmousa, J., Jahed, H., and S. Lambert. 2011.Cyclic axial and cyclic torsional behaviour of extrudedAZ31B magnesium alloy. Int. J. Fatigue 33:1403–16.

122. Dieter, G. E. 1986. Mechanical Metallurgy. 3rd Ed.,

McGraw-Hill, Boston, Massachusetts. pp. 197–01.

123. Yoo, M. H., Morris, J. R., Ho, K. M., and S. R. Agnew.2002. Nonbasal deformation modes of HCP metals and alloys:Role of dislocation source and mobility. Metall. Mat.Trans. A. 33A:813–22.

124. Serra, A., Bacon, D. J., and R. C. Pond. 2002. Twinsas barriers to basal slip in hexagonal close-packedmetals. Metall. Mat. Trans. A. 33A:809–12.

125. Jiang, L., Jonas, J. J., Luo, A. A., Sachdev, A. K.,and S. Godet. 2006. Microstructural and texture evolutionduring the uniaxial tensile testing of AM30 magnesiumalloy. Magnesium Tech. TMS, 135th Annual Meeting. 233–38.

126. Xu, S., Gertsman, V. Y., Li, J., Thompson, J. P., andM. Sahoo. 2005. Role of mechanical twinning in tensilecompressive yield asymmetry of die cast Mg alloys. CanadianMetall. Quart. 44(2):155–66.

127. Jain, J., Poole, W. J., Sinclair, C. W., and M. A.,Gharghouri. 2010. Reducing the tension-compression yieldasymmetry in a Mg-8Al-0.5Zn alloy via precipitation.Scripta Mater. 62(5):301–04.

128. Chen, L. J., Wang, C. Y., Wu, W., Liu, Z., Stoica, G.M., Wu, L., and K. Liaw. 2007. Low-cycle fatigue behaviorof an As-extruded AM50 magnesium alloy. Metall. Mater.Trans. A. 38:2235–41.

129. Chai, H. F., and C. Laird. 1987. Mechanisms of cyclicsoftening and cyclic creep in low carbon steel. Mater.Sci. Eng. 93:159–74.

130. Laird, C. 1981. Cyclic deformation, fatigue cracknucleation and propagation in metals and alloys. InMetallurgical Treatises. Edited by J. K. Tien, and J. F.Elliott. The Metallurgical Society of the AIME,Warrendale, PA, pp. 505–28.

131. Reed-Hill, R. E., and R. Abbaschinan. 1994. PhysicalMetallurgy Principles, 3rd ed. PWS Publishing Company,Boston, MA.

132. Park, S. H., Hong, S., and C.S. Lee. 2010. Activationmode dependent {10–12} twinning characteristics in apolycrystalline magnesium alloy. Scripta Mater. 62:202–05.

133. Kwon, S. H., Song, K. S., Shin, K. S., and S. I. Kwun.

2011. Low cycle fatigue properties and an energybasedapproach for as-extruded AZ31 magnesium Alloy. Met. Mater.Int. 17(2):207–13.

134. Yoo, M. H., Morris, J. R., Ho, K. M., and S. R. Agnew.2001. Non-basal slip systems in HCP metals and alloys:Source mechanisms. Mater. Sci. Eng. A. A319-A321:87–92.

135. Lou, X.Y., Li, M., Boger, R. K., Agnew, S. R., and R.H. Wagoner. 2007. Hardening evolution of AZ31B Mg sheet.Int. J. Plast. 23:44–86.

136. Brown, D. W., Agnew, S. R., Abeln, S. P., Blumenthal,W. R., Bourke, M. A. M., Mataya, M. C., Tomé, C. N., andS. C. Vogel. 2005. The role of texture, temperature, andstrain rate in the activity of deformation twinning. Mat.Sci. Forum 495–497:1037–42.

137. Gharghouri, M. A., Weatherly, G. C., Embury, J. D.,and J. Root. 1999. Study of the mechanical properties ofMg-7.7at.% Al by in-situ neutron diffraction. Philos. Mag.A. 79:1671–95.

138. Oliver, E. C., Daymond, M. R., and P. J. Withers.2005. Neutron diffraction study of extruded magnesiumduring cyclic and elevated temperature loading. Mater. Sci.Forum 490–491:257–62.

139. Patel, H. A., Chen, D. L., Bhole, S. D., and K.Sadayappan. 2010. Cyclic deformation and twinning in asemi-solid processed AZ91D magnesium alloy. Mater. Sci.Eng. A. 528(1):208–19.

140. Jordon, J. B., Gibson, J. B., Horstemeyer, M. F.,Kadiri, H. El., Baird, J. C., and A. A. Luo. 2011. Effectof twinning, slip, and inclusions on the fatigue anisotropyof extrusion-textured AZ61 magnesium alloy. Mater. Sci.Eng. A. 528:6860–71.

141. Yin, S. M., Yang, H. J., Li, S. X., Wu, S. D., and F.Yang. 2008. Cyclic deformation behavior of asextrudedMg-3%Al-1%Zn. Scripta Mater. 58(9):751–54.

142. Yin, S. M., Yang, F., Yang, X. M., Wu, S. D., Li, S.X., and G. Y., Li. 2008. The role of twinning-detwinning onfatigue fracture morphology of Mg-3%Al-1%Zn alloy. Mater.Sci. Eng. A. 494(1–2):397–00.

143. Yu, Q., Zhang, J. X., Jiang,Y. Y., and Q. Z. Li. 2012.An experimental study on cyclic deformation and fatigue of

extruded ZK60 magnesium alloy. Int. J. Fatigue. 36:47–58.

144. Huppmann, M., Lentz, M., Chedid, S., and R. Walter.2011. Analyses of deformation twinning in the extrudedmagnesium alloy AZ31 after compressive and cyclic loading.J. Mater. Sci. 46(4):938–50.

145. Park, S. H., Hong, S. G., Byoung, H. L., Wonkyu, B.,and Chong, S. L. 2010. Low-cycle fatigue characteristics ofrolled Mg-3Al-1Zn alloy. Int. J. Fatigue 32(11):1835–42.

146. Wu, L., Agnew, S. R., Ren,Y., Brown, D. W., Clausen,B., Stoica, G. M., Wenk, H. R., and P. K. Liaw. 2010. Theeffects of texture and extension twinning on the low-cyclefatigue behavior of a rolled magnesium alloy AZ31B. Mater.Sci. Eng. A. 527(26):7057–67.

147. Luo, T. J., Yang, Y. S., Tong, W. H., Duan, Q. Q., andX. G. Dong. 2010. Fatigue deformation characteristic ofas-extruded AM30 magnesium alloy. Mater. Design 31:1617–21.

148. Lin, X. Z., and D. L. Chen. 2008. Strain hardening andstrain-rate sensitivity of an extruded magnesium alloy. J.Mater. Eng. Per. 17:894–01.

149. Decker, R. F. The renaissance in magnesium. 1998. Adv.Mater. Process. 154:31–33.

150. Li, F., Wang, Y., Chen, L., Liu, Z., and J. Zhou.2005. Low-cycle fatigue behavior of two magnesium alloys.J. Mater. Sci. 40:1529–31. For catalo ing purposes only

151. Agnew, S. R., and O. Duygulu. 2005. Plastic anisotropyand the role of non-basal slip in magnesium alloy AZ31B.Int. J. Plas. 21:1161–93.

152. Koike. J. 2005. Enhanced deformation mechanisms byanisotropic plasticity in polycrystalline Mg alloys atroom temperature. Metall. Mater. Trans. A. 36:1689–96.

153. Gerard, B. F. 2010. Anisotropy and voiding at highstrain rates in a Mg alloy extrudate. Adv. Mater. Process.168:32–33.

154. Lv, F., Yang, F., Duan, Q. Q., Yang, Y. S., Wu, S. D.,Li, S. X., and Z. F. Zhang. 2011. Fatigue properties ofrolled magnesium alloy (AZ31) sheet: In�uence of specimenorientation. Int. J. Fatigue 33:672–82.

155. Jiang, L., Jonas, J. J., Mishra, R. K., Luo, A. A.,

Sachdev, A. K., and S. Godet. 2007. Twinning and texturedevelopment in two Mg alloys subjected to loading alongthree different strain paths. Acta Mater. 55:3899–10.

156. Brown, D. W., Agnew, S. R., Bourke, M. A. M., Holden,T. M., Vogel, S. C., and C. N. Tome. 2005. Internal strainand texture evolution during deformation twinning inmagnesium. Mater. Sci. Eng. A. 399:1–12.

157. Grinberg, N. M., Serdyuk, V. A., Yakovenko, L. F.,Malinkina, T. I., and A. S. Kamyshkov. 1977. Kinetics andmechanism of fatigue fracture of magnesium alloys MA2–1 andMA12. Probl. Prochn. 8:40–45.

158. Serdyuk, V. A., and N. M. Grinberg. 1980. Resofteningof IMV6 magnesium alloy in the fatigue process. Probl.Prochn. 1:35–39.

159. Bhambri, A. K., and T. Z. Kattamis. 1971. Castmicrostructure and fatigue behaviour of a grain-re�nedMg-Zn-Zr alloy. Met. Trans. 2:1869–74.

160. Gregory, J. K. 1987. Proceedings of the 3rdInternational Conference of Fatigue Fatigue Thresholds.Charlottesville, pp. 303–13.

161. Horstemeyer, M. F., Oglesby, D., and Fan, J. 2006.CAVS/MSU Report Prepared for USCAR/USAMPAMD, Center forAdvanced Vehicular Systems (CAVS), Starkville, MS.

162. Mayer, H., Papakyriacou, M., Zettl, B., and S. E.Stanzl-Tschegg. 2003. In�uence of porosity on the fatiguelimit of die cast magnesium and aluminium alloys. Int. J.Fatigue 25:245–56.

163. Laird. C. 1967. Fatigue Crack Propagation, ASTM, WestConshohocken, PA. ASTM STP 415.

164. Mansoor, B., Mukherjee, S., and A. Ghosh. 2009.Microstructure and porosity in thixomolded Mg alloys andminimizing adverse effects on formability. Mater. Sci. Eng.A. 512:10–18.

165. D’Errico, F., Rivolta, B., Gerosa, R., and G.Perricone. 2008. Thixomolded magnesium alloys: Strategicproduct innovation in automobiles. JOM 60:70–76.

166. Xu, Y. L., Zhang, K., Li, X. G., Lei, E. J., Yang Y.S., and T., and J. Luo. 2008. High cycle fatigue propertiesof die-cast magnesium alloy AZ91D-1%MM. Trans. Nonferrous

Met. Soc. China 18:s306-s11.

167. Moore, A. R., Torbet, C. J., Shyam, A., Jones, J. W.,Walukas, D. M., and R. F. Decker. 2004. Fatigue Behaviorof Thixomolded ® Magnesium AZ91D Using UltrasonicTechniques. Magnesium Technology 2004. The Minerals,Metals & Materials Society, United States, pp. 263–68.

168. Yu, Q., Zhang, J., Jiang, Y., and Q. Li. 2011.Multiaxial fatigue of extruded AZ61A magnesium alloy. Int.J. Fatigue 33(3):437–47.

169. Albinmousa, J., Jahed, H., and S. Lambert. 2011.Cyclic behaviour of wrought magnesium alloy undermultiaxial load. Int. J. Fatigue 33:1127–39.

170. Liu, L., and C. Dong. 2006. Gas tungsten-arc �llerwelding of AZ31 magnesium alloy. Mater. Lett. 60:2194–97.

171. Chowdhury, S. M., Chen, D. L., Bhole, S. D., and X.Cao. 2010. Effect of pin tool thread orientation onfatigue strength of friction stir welded AZ31B-H24 Mg buttjoints. Proc. Eng. 2(1):825–33.

172. Chan, K. S., Yi-Ming, P., Davidson, D., and R. C.McClung. 1997. Fatigue crack growth mechanisms in HSLA-80steels. Mater. Sci. Eng. A. A222:1–8.

173. Chen, J., Wang, J. Q., Han, E. H., and W. Ke. 2008.Electrochemical corrosion and mechanical behaviors of thecharged magnesium. Mater. Sci. Eng. A. 494(1–2):257–62.

174. Chen, J., Wang, J. Q., Han, E. H., Dong, J. H., and W.Ke. 2007. AC impedance spectroscopy study of the corrosionbehavior of an AZ91 magnesium alloy in 0.1 M sodium sulfatesolution. Electrochim. Acta 52(9):3299–09.

175. Zhou, W. Q., Shan, D. Y., Han, E. H., and W. Ke. 2008.Structure and formation mechanism of phosphate conversioncoating on die-cast AZ91D magnesium alloy. Corrosion Sci.50(2):329–37.

176. Winzer, N., Atrens, A., Song, G. L., Ghali, E.,Dietzel, W., Kainer, K. U., Hort, N., and C., Blawert.2005. A critical review of the Stress Corrosion Cracking(SCC) of magnesium alloys. Adv. Eng. Mater. 7(8):659–93.

177. Blawert, C., Dietzel, W., Ghali, E., and G. L., Song.2006. Anodizing treatments for magnesium alloys and theireffect on corrosion resistance in various environments,

Adv. Eng. Mater. 8(6):511–33.

178. Atrens, A., Winzer, N., and W., Dietzel. 2011. Stresscorrosion cracking of magnesium alloys. Adv. Eng. Mater.13(1–2):11–18.

179. Arrabal, R., Matykina, E., Viejo, F., Skeldon, P., andG. E. Thompson. 2008. Corrosion resistance of WE43 andAZ91D magnesium alloys with phosphate PEO coatings.Corrosion Sci. 50(6):1744–52.

180. Huo, H. W., Li, Y., and F. H., Wang. 2004. Corrosionof AZ91D magnesium alloy with a chemical conversion coatingand electroless nickel layer. Corrosion Sci. 46(6):1467–77.

181. Zhang, T., Meng, G. Z., Shao, Y. W., Cui, Z. Y., andF. H., Wang. 2011. Corrosion of hot extrusion AZ91magnesium alloy. Part II: Effect of rare earth elementneodymium (Nd) on the corrosion behavior of extrudedalloy. Corrosion Sci. 53(9):2934–42.

182. Zhang, T., Shao, Y. W., Meng, G. Z., Cui, Z. Y., andF. H., Wang. 2011. Corrosion of hot extrusion AZ91magnesium alloy: I-relation between the microstructure andcorrosion behavior. Corrosion Sci. 53(5):1960–68.

183. Jackman, J. A., Wood, J., Essadiqi, E., Lo, J., Sahoo,M., Xu, S., Thomson, J., and W. J., Liu. 2005. Overview ofkey R&D activities for the development of high-performancemagnesium materials in Canada. Mater. Sci. Forum488–489:21–24.

184. Villafuerte, J., and W. Y., Zheng. 2007. Corrosionprotection of magnesium alloys by cold spray. Adv. Mater.Proc. 165(9):53–54.

185. Hussein, R. O., Zhang, P., Northwood, D. O., and X.,Nie. 2011. Improving the corrosion resistance of magnesiumalloy AJ62 by a plasma electrolytic oxidation PEO coatingprocess. Corrosion Mater. 36(3):38–49.

186. Luo, A., and M. O. Pekguleryuz. 1994. Cast magnesiumalloys for elevated temperature applications. J. Mater.Sci. 29(20):5259–71.

187. Pekguleryuz, M., and M., Celikin. 2010. Creepresistance in magnesium alloys. Int. Mater. Rev.55(4):197–17.

188. Hu, H., Yu, A., Li, N. Y., and J. E., Allison. 2003.

Potential magnesium alloys for high temperature die castautomotive applications: A review. Mater. Manuf.18(5):687–17.

189. Vespa, G., Mackenzie, L. W. F., Verma, R., Zarandi,F., Essadiqi, E., and S., Yue. 2008. The in�uence of theas-hot rolled microstructure on the elevated temperaturemechanical properties of magnesium AZ31 sheet. Mater. Sci.Eng. A. 487(1–2):243–50.

190. Shang, L., Yue, S., Verma, R., Krajewski, P., Galvani,C., and E., Essadiqi. 2011. Effect of microalloying (Ca,Sr, and Ce) on elevated temperature tensile behavior ofAZ31 magnesium sheet alloy. Mater. Sci. Eng. A.528(10–11):3761–70.

191. Werdin, S., Tromann, T., Gugau, M., and K. L. Kotte.2005. Structural durability of magnesium alloys andcomponents. Materialwissenschaft und Werkstofftechnik.36(11):659–68.

192. Hilpert, M., and L. Wagner. 2000. Effect of mechanicalsurface treatment and environment on fatigue behavior ofwrought magnesium alloys. In Magnesium Alloys and TheirApplications, Edited by K. U. Kainer. Wiley-Vch: Weinheim,Germany, pp. 463–68.

193. Gutman, E.M., Unigovski, Ya., Eliezer, A., Abramov,E., and L. Riber. 2000. Effect of processing andenvironment on mechanical properties of die cast magnesiumalloys. Light Metal Age 58:14–20.

194. Zhou, H. M., Wang, J. Q., Zang, Q. S., and E. H. Han.2007. Study on the effect of Cl − concentration on thecorrosion fatigue damage in a rolled AZ31B magnesium alloyby acoustic emission. Key Eng. Mater. 353–358 (PART1):327–30.

195. Eliezer, A., Haddad, J., Unigovski, Y., and E. M.Gutman. 2005. Static and dynamic corrosion fatigue of Mgalloys used in automotive industry. Mater. Manuf. Process.20(1):75–88.

196. Stephens, R. I., Schrader, C. D., and K. B. Lease.1995. Corrosion fatigue of AZ91E-T6 cast magnesiun alloyin a 3.5 percent NaCl aqueous environment. J. Eng. Mater.Tech. 117:293–98.

197. Xue, Y., Horstemeyer, M. F., McDowell, D. L., ElKadiri, H., and J. Fan. 2007. Microstructure-based

multistage fatigue modeling of a cast AE44 magnesium alloy.Int. J. Fatigue 29(4):666–76.

198. Lamark, T. T., Janecek, M., and Y. Estrin. 2005.In�uence of surface skin on the fatigue properties ofdiecast magnesium alloy AS21X. Mater. Sci. Forum482:379–82.

199. Mayer, H., Papakyriacou, M., Zettl, B., and S. Vacic.2005. Endurance limit and threshold stress intensity ofdie cast magnesium and aluminium alloys at elevatedtemperatures. Int. J. Fatigue 27(9):1076–88. For at logingpurposes nly

14 Chapter 14 - Aerogel Materialsfor Aerospace

1. Aerogel—Mystifying Blue Smoke.Stardust.jpl.nasa.gov/aerogel_factsheet.pdf. (accessedDecember 21, 2010).

2. Schaedler1, T. A., Jacobsen, A. J., Torrents, A. et al.Ultralight metallic microlattices. Science 334: 962–965.

3. Press Release. Guinness Records Names JPL’s AerogelWorld’s Lightest Solid. 2002. http://stardust.jpl.nasa.gov/news/news93.html. (accessed December 21, 2010).

4. Kistler, S. S. 1931. Coherent expanded aerogels andjellies. Nature 127: 741–741.

5. Kistler, S. S. 1922. Two Component Systems Near theCritical Temperature of One Component. Engineer thesis.Leland Stanford Junior University.

6. Caldwell, A. G. 1934. Thermal conductivity of silicaaerogel. Ind. Eng. Chem. 26: 658–662.

7. Swann, S., Appel, E. G., and Kistler, S. S. 1934. Thoriaaërogel catalyst: Aliphatic esters to ketones. With Appel.Ind. Eng. Chem. 26: 1014–1014.

8. Kistler, S. S. 1935. The relation between heatconductivity and structure in silica aerogel. J. Phys.Chem. 39: 79–86.

9. Advertisement. 1943. Skeleton sand powder tops all heatinsulators. Popular Sci. 142: 80–81.

10. Advertisement. 1954. SANTOCEL. . . Can this remarkablesilica aerogel improve YOUR processing or products? Chem.Eng. News 32: 4179–4179.

11. Advertisement. 1964. A little Santocel ® solves a lotof problems. Chem. Eng. News 42: 55–55.

12. Nicoloan, G. A. 1968. Contribution on l’etude desaerogels des silica. PhD dissertation. University of Lyon.

13. Teichner, S. J., and Nicoloan, G. A. 1972. Method ofPreparing Inorganic Aerogels. United States Patent No.3,672,833.

14. Poelz, G., and Riethmueller R. 1982. Preparation of

silica aerogel for Cerenkov counters. Nuc. Instr. Meth.195: 491–503.

15. Henning, S., and Svensson, L. 1981. Production ofsilica aerogel. Phys. Scr. 23: 697–702.

16. Tewari, P. H., and Hunt, A. J. 1986. Process forForming Transparent Aerogel Insulating Arrays. UnitedStates Patent No. 4,610,863.

17. Tewari, P. H., Hunt, A. J., and Lofftus, K. D. 1985.Ambient-temperature supercritical drying of transparentsilica aerogels. Mater. Lett. 3: 363–367.

18. Herrmann, G., Iden, R., Mielke, M., Teich, F., andZiegler, B. 1995. On the way to commercial production ofsilica aerogel. J. Non-Cryst. Solids 186: 380–387.

19. Pekala, R. W. 1989. Organic aerogels from thepolycondensation of resorcinol with formaldehyde.J. Mater. Sci. 24: 3221–3227.

20. Pekala, R. W., Alviso, C. T., and LeMay, J. D. 1990.Organic aerogels: Microstructural dependence of mechanicalproperties in compression. J. Non-Cryst. Solids 125: 67–75.

21. Pekala, R. W., Mayer, S. T., Kaschmitter J. L., andKong, F. M. 1994. Carbon aerogels: An update on structure,properties, and applications. In Sol–Gel Science Processingand Applications, ed. Y. A. Attia, pp. 369–337. New York:Plenum.

22. Mayer, S. T., Pekala, R. W., and Kaschmitter, J. L.1993. First demonstration of a supercapacitor using carbonaerogel the aerocapacitor: An electrochemical double-layerenergy-storage device. J. Electrochem. Soc. 140: 446–451.

23. Tillotson, T. M., and Hrubesh, L. W., 1992. Transparentultralow-density silica aerogels prepared by a two-stepsol-gel process. J. Non-Cryst. Solids 145: 44–50.

24. Tsou, P. 1995. Aerogel captures cosmic dust intact. J.Non-Cryst. Solids 186: 415–427.

25. Poco, F., Coronado, P. R., Pekala, R. W., and Hrubesh,L. W. 1996. A rapid supercritical extraction process forthe production of silica aerogels. In Materials ResearchSociety Symposium Proceedings, eds. Lobo, R. F., Beck, J.S., Suib, S. L., Corbin, D. R., Davis, M. E., Iton, L. E.,and Zones, S. I., pp. 297–302. Cambridge: Cambridge

University. Press.

26. Poco, J. F., Satcher, Jr. J. H., and Hrubesh, L. W.2001. Synthesis of high porosity, monolithic aluminaaerogels. J. Non-Cryst. Solids 285: 56–63.

27. Fricke, J., and Tillotson T. 1997. Aerogels:Production, characterization, and applications. Thin SolidFilms 297: 212–223.

28. Smith, D. M., Stein, D., Anderson J. M., and Ackerman,W. 1995. Preparation of low-density xerogels at ambientpressure. J. Non-Cryst. Solids 186: 104–112.

29. Scherer, G. W. 1994. Relaxation of a viscoelastic gelbar: II. Silica gel. J. Sol–Gel Sci. Techno. 3: 199–204.

30. Prakash, S. S., Brinker, C. J., and Hurd, A. J. 1995.Silica aerogel �lms at ambient pressure. J. Non-Cryst.Solids 190: 264–275.

31. Prakash, S. S., Brinker, C. J., Hurd, A. J., and Rao,S. M. 1994. Silica aerogel �lms prepared at ambientpressure by using surface derivatization to inducereversible drying shrinkage. Nature 374: 439–443.

32. Tillotson, T. M., Sunderland, W. E., Thomas, I. M., andHrubesh, L. W. 1994. Reintroduction of epoxideassistedgelation for preparing metal oxide aerogels synthesis oflanthanide and lanthanide-silicate aerogels. J. Sol–GelSci. Techno. 1: 241–249.

33. Burnett, D. S. 2006. NASA returns rocks from a Comet.Science 314: 1709–1710.

34. Leventis, N., Sotiriou-Leventis, C., Zhang, G., andAbdel-Monem M. R. 2002. Nanoengineering strong silicaaerogels. Nano Lett. 2: 957–960.

35. Leventis, L., Mulik, S., and Sotiriou-Leventis, C.2007. Crosslinking 3d assemblies of silica nanoparticles(aerogels) by surface-initiated free radical polymerizationof styrene and methylmethacrylate. Polymer Preprints 48:950–951.

36. Leventis, N., Chandrasekaran, N. Sotiriou-Leventis, C.,and Mumtaz, A. 2009. Smelting in the age of nano: Ironaerogels. J. Mater. Chem. 19: 63–65.

37. Mohanan, J. L., Arachchige, I. U., and Brock, S. L.

2005. Synthetic pathway for producing semiconducting metalchalcogenide aerogels porous semiconductor chalcogenideaerogels. Science 307: 397–401.

38. Kistler, S. S. 1932. Coherent expanded-aerogels. J.Phys. Chem. 34: 52–64.

39. Dorcheh A., and Abbasi, M. H. 2008. Silica aerogel:Synthesis, properties and characterization. J. Mater.Process. Tech. 199: 10–26.

40. Moner-Girona, M., Roig, A., and Molins, E. 2003.Sol-gel route to direct formation of silica aerogelmicroparticles using supercritical solvents. J. Sol–GelSci. Tech. 26: 645–649.

41. Reichenauer, G. 2004. Thermal aging of silica gels inwater. J. Non-Cryst. Solids 350: 189–195.

42. Einasrud, M. A., Dahle, M., Lima, S., and Hæreid, S.1995. Preparation and properties of monolithic silicaxerogels from TEOS-based alcogels aged in silane solutions.J. Non-Cryst. Solids 186: 96–103.

43. Hæreid, S., Anderson, J., and Einarsrud, M. A. 1995.Thermal and temporal aging of TMOS-based aerogel precursorsin water. J. Non-Cryst. Solids 185: 221–226.

44. Haereid, S., Nilsen, E., Ranum, V., and Einarsrud,M.-A. 1997. Thermal and temporal aging of two stepsacid-base catalyzed silica gels in water/ethanol solutions.J. Sol-Gel Sci. Tech. 8:153–157.

45. Einarsrud, M. A., and Nilsen, E. 1998. Strengthening ofwater glass and colloidal sol based silica gels by agingin TEOS. J. Non-Cryst. Solids 226: 122–128.

46. Smitha, S., Shajesh, P., Aravind, P. R., Kumar, S. R.,Pillai, P. K., and Warrier, K. G. K. 2006. Effect of agingtime and concentration of aging solution on the porositycharacteristics of subcritically dried silica aerogels.Microporous Mesoporous Mater. 91: 286–292.

47. Pajonka, G. M., Raob, A. V., Sawantc, B. M., andParvathy, N. N. 1998. Dependence of monolithicity andphysical properties of TMOS silica aerogels on gel agingand drying conditions. J. Non-Cryst. Solids 209: 40–50.

48. Hdacha, H., Woigniera, T., Phalippoua, J., andSchererb, G. W. 2003. Effect of aging and pH on the

modulus of aerogels. J. Non-Cryst. Solids 121: 202–205.

49. Keysar, S., Shter, G. E., de Hazan, Y., Cohen, Y., andGrader, G. S. 1997. Direct formation of aerogels bySol−Gel polymerizations of alkoxysilanes in supercriticalcarbon dioxide. Chem. Mater. 9: 2464–2467.

50. Yokogawa, H., and Yokoyama, M. 1995. Hydrophobic silicaaerogels. J. Non-Cryst. Solids 186: 23–29.

51. Štandekera, S., Novaka, Z., and Knez, Ž. 2007.Adsorption of toxic organic compounds from water withhydrophobic silica aerogels. J. Colloid Interface Sci. 310:362–368.

52. Reynolds, J. G., Coronado, P. R., and Hrubesh, L. W.2001. Hydrophobic aerogels for oil-spill cleanup—Synthesis and characterization. J. Non-Cryst. Solids292: 127–137.

53. Scherer, G. W. 1993. Freezing gels. J. Non-Cryst.Solids 155: 1–25.

54. Bryning, M. B., Milkie, D. E., Islam, M. F., Hough, L.A., Kikkawa, J. M., and Yodh, A. G. 2007. Carbon nanotubeaerogels. Adv. Mater. 19: 661–664.

55. Casas, L., Roig, A., Rodríguez, E., Molins, E., Tejada,J., and Sort J. 2001. Silica aerogel–iron oxidenanocomposites: Structural and magnetic properties. J.Non-Cryst. Solids 285: 37–43.

56. Rodembusch, F. S., Campo, L. F., Stefani, V., andRigacci, A. 2005. The �rst silica aerogel �uorescent byexcited state intramolecular proton transfer mechanism(ESIPT). J. Mater. Chem. 15:1537–1541.

57. Bockhorst, M., Heinlotha, K., Pajonkb, G. M., Begagb,R., and Elaloui, E. 1995. Fluorescent dye doped aerogelsfor the enhancement of Čerenkov light detection. J.Non-Cryst. Solids 186: 388–232.

58. Ayers, M. R., and Hunt, A. J. 1998. Molecular oxygensensors based on photoluminescent silica aerogels. J.Non-Cryst. Solids 225: 343–347.

59. Hunt, A. J., Ayers, M. R., and Cao, W. 1995. Aerogelcomposites using chemical vapor in�ltration. J. Non-Cryst.Solids 185: 227–232. F r cataloging purposes ly

60. Bednarczyk, S., Bechir, R., and Baclet, P. 1999. Lasermicromachining of small objects for high energy laserexperiments. Appl. Phys. A 69: 495–500.

61. Prassas, M., Phalippou, J., and Zarzycki, J. 1986.Sintering of monolithic silica aerogels. In Science ofCeramic Chemical Processing, eds. Hench, L. L., and Ulrich,D. R., pp. 156–167. New York: Wiley.

62. Kim, J. H., Suh, D. J., Park, T. J., and Kim, K. L.2000. Effect of metal particle size on coking during CO 2reforming of CH 4 over Ni–alumina aerogel catalysts. Appl.Catal. A-Gen. 197: 191–200.

63. Bono Jr. M. S., Anderson, A. M., and Carroll, M. K.2010. Alumina aerogels prepared via rapid supercriticalextraction. J. Sol-Gel Sci. Technol. 53: 216–226.

64. Campbell, L. K., Na, B. K., and Ko, E. I. 1992.Synthesis and characterization of titania aerogels. Chem.Mater. 4: 1329–1333.

65. Ward, D. A., and Ko, E. I. 1993. Synthesis andstructural transformation of zirconia aerogels. Chem.Mater. 5: 956–969.

66. Maurer, S. M., and Ko, E. I. 1992. Structural andacidic characterization of Niobia aerogels. J. Catal.135:125–134.

67. Livage, J., Henry, M., and Sanchez, C. 1988. Sol-gelchemistry of transition metal oxides. Prog. Solid StateChem. 18: 259.

68. Yoldas, B. E. 1975. Alumina gels that form poroustransparent Al 2 O 3 . J. Mater. Sci. 10: 1856–1860.

69. Gash, A. E., Tillotson, T. M., Satcher Jr. J. H.,Hrubesh, L. W., and Simpson, R. L. 2001. New sol–gelsynthetic route to transition and main-group metal oxideaerogels using inorganic salt precursors. J. NonCryst.Solids 285: 22–28.

70. Gao, Y., Sisk C., and Hope-Weeks, L. J. 2007. A sol-gelroute to synthesize monolithic zinc oxide aerogels. Chem.Mater. 19: 6007–6011.

71. Rigacci A., and Achard P. 2011. Cellulosic andPolyurethane Aerogels. In Aerogels Handbook, ed. M. A.Aegerter, N. Leventis, and M. M. Koebel, pp.191–214. New

York: Springer-Verlag.

72. Kim, S. Y., Yeo, D. H., Lim, J. W., Yoo, K., Lee, K.,and Kim, H. 2001. Synthesis and characterization ofresorcinol-formaldehyde organic aerogel. J. Chem. Eng. Jpn.34: 216–220.

73. Pekala, R. W., and Alviso, C. T. 1990. A new syntheticroute to organic aerogels. Mater. Res. Soc. Symp. 180:791–796.

74. Saliger, R., Fischer, U., Herta C., and Fricke J. 1998.High surface area carbon aerogels for supercapacitors. J.Non-Cryst. Solids 225: 81–85.

75. Long, J. W., Dunn, B., Rolison, D. R., and White, H. S.2004. Three-dimensional battery architectures. Chem. Rev.104: 4463–4492.

76. Li, G. R., Feng, Z.-P., Ou, Y.-N., Wu, D., Fu, R., andTong, Y.-X. 2010. Mesoporous MnO 2 /carbon aerogelcomposites as promising electrode materials forhigh-performance supercapacitors. Langmuir 26: 2209–2213.

77. Kalpana, D., Omkumar, K., Kumar, S., and Renganathan N.2006. A novel high power symmetric ZnO/carbon aerogelcomposite electrode for electrochemical supercapacitor.Electrochim. Acta 52: 1309–1315.

78. Pröbstle, H., Schmitt, C., and Frick, J. 2002. Buttoncell supercapacitors with monolithic carbon aerogels. J.Power Sources 105: 87–92.

79. Xu, P., Drewes, J. E., Heil, S., and Wang, G. 2008.Treatment of brackish produced water using carbonaerogel-based capacitive deionization technology. Water.Res. 42: 2605–2617.

80. Mechanically Strong, Lightweight Porous MaterialsDeveloped (X-Aerogels). http://www.grc.nasa.gov/WWW/RT/2004/RM/RM11P-leventis.html.

81. Meador, M. A. B., Johnston, J. C., Fabrizio, E. F.,Ulvi F., and Leventis, N. 2010. Highly Porous andMechanically Strong Ceramic Oxide Aerogels. United StatesPatent No. 7,732,496.

82. Capadona, L. A., Meador, M. A. B., Alunni, A.,Fabrizio, E. F., Vassilaras, P., and Leventis, N. 2006.Flexible, low-density polymer crosslinked silica aerogels.

Polymer 47: 5754–5761.

83. Katti, A., Shimpi, N., Roy, S. et al. 2006. Chemical,physical, and mechanical characterization of isocyanatecross-linked amine-modi�ed silica aerogels. Chem. Mater.18: 285–296.

84. Mulik, S., Sotiriou-Leventis, C., and Leventis, N.2008. Macroporous electrically conducting carbon networksby pyrolysis of isocyanate-cross-linkedresorcinol-formaldehyde aerogels. Chem. Mater. 20:6985–6997.

85. Nyce, G. W., Hayes, J. R., Hamza, A. V., and SatcherJr. J. H. 2007. Synthesis and characterization ofhierarchical porous gold materials. Chem. Mater. 19:344–346.

86. Tappan, B., Steiner, S., and Luther, E. 2010.Nanoporous metal foams. Angew. Chem. Int. Ed. 49:4544–4565.

87. Bag, S., Trikalitis, P. N., Chupas, P. J., Armatas, G.S., and Kanatzidis, M. G. 2007. Porous semiconducting gelsand aerogels from chalcogenide clusters. Science 317:490–493.

88. Arachchige, I., and Brock, S. 2007. Sol-gel methods forthe assembly of metal chalcogenide quantum dots. Acc.Chem. Res. 40: 801–809.

89. Zheng, B., Li, Y., and Liu, J. 2002. CVD synthesis andpuri�cation of single-walled carbon nanotubes onaerogel-supported catalyst. App. Phys. A 74: 345–348.

90. Fu. R., Dresselhaus, M. S., Dresselhaus, G. et al.2003. The growth of carbon nanostructures on cobaltdopedcarbon aerogels. J. Non-Cryst. Solids 318: 223–232.

91. Bordjiba, T., Mohamedi, M., and Dao, L. H. 2008. Newclass of carbon-nanotube aerogel electrodes forelectrochemical power sources. Adv. Mater. 20: 815–819.

92. Li, Y.-L., Kinloch, I. A., and Windle, A. H. 2004.Direct spinning of carbon nanotube �bers from chemicalvapor deposition synthesis. Science 304: 276–278.

93. Zhang, M., Fang, S. L., Zakhidov, A. A. et al. 2005.Strong, transparent, multifunctional, carbon nanotubesheets. Science 309: 1215–1219.

94. Aliev, A. E., Oh, J., Kozlov, M. E., Kuznetsov, A. A.,Fang, S., Fonseca, A. F., Ovalle, R. et al. 2009.Giant-stroke, superelastic carbon nanotube aerogel muscles.Science 323: 1575 Article CAS.

95. Zou, J., Liu, J., Karakoti, A. H., Kumar, A. et al.2010. Ultralight multiwalled carbon nanotube aerogel. ACSNano 4: 7293–7302.

96. Worsley, M. A., Pauzauskie, P. J., Olson, T. Y.,Biener, J., Satcher, Jr. J. H., and Baumann, T. F., 2010.Synthesis of graphene aerogel with high electricalconductivity. J. Am. Chem. Soc. 132: 14067–14069.

97. Zhang, X., Sui, Z., Xu, B., Yue, S., Luo,Y., Zhan, W.,and Liu, B. 2011. Mechanically strong and highlyconductive graphene aerogel and its use as electrodes forelectrochemical power sources. J. Mater. Chem. 21:6494–6497.

98. Bisson, A., Rigacci, A., Lecomte, D., Rodier, E., andAchard, P. 2003. Drying of silica gels to obtain aerogels:Phenomenology and basic techniques. Drying Techno. 21:593–628.

99. Morris, P. M., and Shrinivasan, S. 2005. Aerogels:Unique material, fascinating properties and unlimitedapplications. Annual Review of Heat Transfer: In Memory ofChang-Lin Tien, Vol. 14, pp. 385–408. New York: BegellHouse Inc.

100. Iler, R. K. 1979. Silica gels and powders—Physicalcharacterization. In The Chemistry of Silica, pp. 462–510.New York: John Wiley and Sons, Inc.

101. Scherer, G. W. 1989. Effect of shrinkage on themodulus of silica gel. J. Non-Cryst. Solids 109: 183–190.

102. Hüsing, N., and Schubert, U. 1998. Aerogels—Airymaterials: Chemistry, structure, and properties. Angew.Chem. Int. Ed. 37: 22–45.

103. McCusker, L. B., Liebau, F., and Engelhardt, G. 2001.Nomenclature of structural and compositionalcharacteristics of ordered microporous and mesoporousmaterials with inorganic hosts. Pure Appl. Chem. 73:381–394.

104. Brunauer, S., Emmett, P. H., and Teller, E. 1938.

Adsorption of gases in multimolecular layers. J. Am. Chem.Soc. 60: 309–319.

105. Rao, A. V., and Wagh, P. B. 1998. Preparation andcharacterization of hydrophobic silica aerogels. Mater.Chem. Phys. 53: 13–18.

106. Kim, G. S., and Hyun, S. H. 2004. Synthesis andcharacterization of silica aerogel �lms for inter-metaldielectrics via ambient drying. Thin Solid Films 460,190–200.

107. Reichenauer, G., and Scherer, G. W. 2001. Nitrogensorption in aerogels. J. Non-Cryst. Solids 285: 167–174.

108. Vacher, R., Woignier, T., Pelous, J., and Courtens, E.1988. Structure and self-similarity of silica aerogels.Phys. Rev. B 37: 6500–6503.

109. Giesche, H. 2006. Mercury porosimetry: A general(practical) overview. Part. Part. Syst. Charact. 23: 1–11.

110. Schultz, P. J., and Lynn, K. G. 1988. Interaction ofpositron beams with surfaces, thin �lms, and interfaces.Rev. Mod. Phys. 60: 701–779.

111. Mincov, I., Petkov, M. P., Tsou, P., and Troev, T.2004. Porosity characterization of aerogels using positronannihilation lifetime spectroscopy. J. Non-Cryst. Solids350: 253–258.

112. Woignier, T., and Phalippou, J. 1987. Skeletal densityof silica aerogels. J. Non-Cryst. Solids 93: 17–21.

113. Hrubesh, L. W., and Pekala, R. W. 1994. Thermalproperties of organic and inorganic aerogels. J. Mater.Res. 9: 731–738.

114. Smith, D., Scherer, G., and Anderson, J. 1995.Shrinkage during drying of silica aerogel. J. Non-Cryst.Solids 188: 191–206. For cataloging purpo es only

115. Woignier, T., Reynes, J., Alaoui, A. H., Beurroies,I., and Phalippou, J. 1998. Different kinds of structuresin aerogels: Relationships with mechanical properties. J.Non-Cryst. Solids 241: 45–52.

116. Parmenter K. E., and Milstein, F. 1998. Mechanicalproperties of silica aerogels. J. Non-Cryst. Solids 223:179–189.

117. Gross, J., Reichenauer, G., and Fricke, J. 1998.Mechanical properties of silica aerogels. J. Phys. D: Appl.Phys. 21: 1447–1451.

118. Stark, R. W., Drobek, T., Weth, M., Fricke, J., andHekel, W. M. 1998. Determination of elastic properties ofsingle aerogel power particles with the AFM.Ultramicroscopy 75: 161–169.

119. Richter, K., Norris, P. M., and Tien, C. L. 1996.Aerogels: Applications, structure, and heat transferphenomena. Ann. Rev. Heat Transfer 6: 61–114.

120. Caps, R., and Fricke, J. 1986. Radiative heat transferin silica aerogel. In Aerogels, Vol. 6, ed. J. Ficke, pp.110–115. Berlin: Springer-Verlag.

121. Lee, D., Stevens, P. C., Zeng, S. Q., and Hunt, A. J.1995. Thermal characterization of carbon-opaci�ed silicaaerogels. J. Non-Cryst. Solids 186: 285–290.

122. Worsley, M. A., Satcher Jr. J. H., and Baumann T. F.2009. Enhanced thermal transport in carbon aerogelnanocomposites containing double-walled carbon nanotubes.J. Appl. Phys. 105, article ID 084316 (April): 1–4.

123. Kuhn, J., Gleissner, T., Aruini-Schuster, M. C.,Korder, S., and Fricke, J. 1995. Integration of mineralpowders into SiO 2 aerogels. J. Non-Cryst. Solids 186:291–295.

124. Zeng, S., Hunt, A., and Greif, R. 1995. Theoreticalmodeling of carbon content to minimize heat transfer insilica aerogel. J. Non-Cryst. Solids 186: 271–277.

125. Haranath, D., Pajonk, G. M., Wagh, P. B., and Rao, A.V. 1997. Effect of sol-gel processing parameters onthermal properties of silica aerogels. Mater: Chem. Phys49: 129–134.

126. Kim G.-S., and Hyun, S. H. 2003. Synthesis of windowglazing coated with silica aerogel �lms via ambient drying.J. Non-Cryst. Solids 320: 125–132.

127. Bock, V., Nilsson, O., Blumm, J., and Fricke, J. 1995.Thermal properties of carbon aerogels. J. NonCryst. Solids185: 233–239.

128. Beck, A., Linsmeier, J., Körner, W., Scheller, H., and

Fricke J. 1995. Radiation transport and image transmissionthrough aerogels. J. Non-Cryst. Solids 186: 232–237.

129. Hunt, A. J. 1998. Light scattering for aerogelcharacterization. J. Non-Cryst. Solids 225: 303–306.

130. Sumiyoshi, T., Adachi, I., Enomoto, R., Iijima, T.,Suda R., Yokoyam, M., and Yokogawa, H. 1998. Silicaaerogels in high energy physics. J. Non-Cryst. Solids 225:369–374.

131. Jensen, K. I., Schultz, J. M., and Kristiansen, F. H.,2004. Development of windows based on highly insulatingaerogel glazings. J. Non-Cryst. Solids 350: 351–357.

132. Perego, D. L. 2008. Ageing tests and recoveryprocedures of silica aerogel. Nucl. Instr. Meth. A 595:224–227.

133. Buzykaev, A. R., Danilyuk, A. F., Ganzhur, S. F.,Kravchenko, E. A., and Onuchin, A. P. 1999. Measurement ofoptical parameters of aerogel. Nucl. Instr. Meth. A 433:396–400.

134. Rao, A. V., Nilsen, E., and Einarsrud, M. A. 2001.Effect of precursors, methylation agents and solvents onthe physicochemical properties of silica aerogels preparedby atmospheric pressure drying method. J. Non-Cryst.Solids 296: 165–171.

135. Danilyuk, A. F., Kravchenko, E. A., Okunev, A. G.,Onuchin, A. P., and Shaurman, S. A. 1999. Synthesis ofaerogel tiles with high light scattering length. Nucl.Instr. Meth. A 433: 406–407.

136. Mohanan, J. L., and Brock, S. L. 2004. A new additionto the aerogel community: Unsupported CdS aerogels withtunable optical properties. J. Non-Cryst. Solids 350: 1–8.

137. Yao, Q., and Brock, S. L. 2010. Optical sensing oftriethylamine using CdSe aerogels. Nanotechnology, Vol.21, 115502: 1–10.http://iopscience.iop.org/0957–4484/21/11/115502.

138. Jo, M.-H., Hong, J.-K., Park, H.-H., Kim, J.-J., Hyun,S.-H., and Choi, S.-Y. 1997. Application of SiO 2 aerogel�lm with low dielectric constant to intermetal dielectrics.Thin Solid Films 308–309: 490–494.

139. Kawakami, N., Fukumoto, Y., Kinoshita, T., Suzuki, K.,

and Inoue, K. 2000. Preparation of highly porous silicaaerogel thin �lm by supercritical drying. Jpn. J. Appl.Phys. 39: L182–L184.

140. Forest, L., Gibiat, V., and Woignier, T. 1998. Biot’stheory of acoustic propagation in porous media applied toaerogels and alcogels. J. Non-Cryst. Solids 225: 287–292.

141. Gibiata, V., Lefeuvre, O., Woignier, T., Pelous, J.,and Phalippou, J. 1995. Acoustic properties and potentialapplications of silica aerogels. J. Non-Cryst. Solids, 186,2 June: 244–255.

142. Burger, T., and Fricke, J. 1998. Aerogels: Production,modi�cation and applications. Phys. Chem. Chem. Phys. 102:1523–1528.

143. Tsou, P. 1995. Silica aerogel captures cosmic dustintact. J. Non-Cryst. Solids 186: 415–427.

144. Burchell. M. J., Thompson R., and Yano H. 1999.Capture of hypervelocity particles in aerogel: In groundlaboratory and low earth orbit. Planet. Space Sci. 47:189–204.

145. Hörz, F., Zolensky, M. E., Berhard, R. P., See, T. H.,and Warren, J. L. 2000. Impact features and projectilesresidues in aerogel exposed on Mir. Icarus 147: 559–579.

146. Kitazawa, Y., Fujiwara, A., Kadono, T., Imagawa, K.,Okada, K., and Uematsu, K. 1999. Hypervelocity impactexperiments on aerogel dust collector. J. Geophys. Res.104: 22035–22052.

147. Okudairi, K., Noguchi, T., Nakamura, T., Sugita, S.,Sekine, Y., and Yano, H. 2004. Evaluation of mineralogicalalteration of micrometeoroid analog materials captured inaerogel. Adv. Space Res. 34: 2299–2304.

148. Jones, S. M. 2010. Non-silica aerogels ashypervelocity particle capture materials. Meteor. Planet.Sci. 45: 91–98.

149. Shrine, N. R. G., McDonnell, J. A. M., Burchell, M. J.et al. 1997. EuroMir ‘95: �rst results from thedustwatch-P detectors of the European space exposurefacility. Adv. Space Res. 20: 1481–1484.

150. Kitazawa, Y., Fujiwara, A., Toshihiko, K. et al.2008. Overview of the MPAC experiment—development of dust

collectors, hypervelocity impact experiments, and post�ight analysis. In Proceedings of International Symposiumon “SM/MPAC&SEED Experiment”, JAXA Special Publication(JAXA-SP08–015E), pp. 49–58. Japan: JAXA.

151. Hörz, F., Cross, G., Zolensky, M. E., See, T. H. E.,Bernhard, R.P., and Warren, J. L. 1999. Optical analysis ofimpact features in aerogel from the orbital debriscollection experiment on the Mir station, NASATM-1999-209372.

152. Ishii, H. A., Graham, G. A., Kearsley, A. T., Grant,P. G., Snead, C. J., and Bradley, J. P., 2005. Ultrasonicmicro-blades for the rapid extraction of impact tracks fromaerogel. In Proc. Lunar Planetary Sci. Conf. 36, abstractno. 1387. Huston. Lunar Planetary Institute.

153. Flynn, G. J., Bleuet, P., Borg, J. et al. 2006.Elemental compositions of Comet 81P/Wild 2 samplescollected by Stardust. Science 316: 1731–1735.

154. McKeegan, K. D., Aléon, J., Bradley, J. et al. 2006.Isotopic compositions of Cometary matter returned byStardust. Science 316: 1724–1728.

155. Sandford, S. A., Aléon, J., Alexander C. M. et al.2006. Organics captured from Comet 81P/Wild 2 by theStardust spacecraft. Science 316: 1720–1724.

156. Elsila, J. E., Glavin, D. P., and Dworkin, J. P. 2009.Cometary glycine detected in samples returned by Stardust.Meteor. Planet. Sci. 44: 1323–1330.

157. Leshin, L. A., Yen, A., Bomba, J. et al. 2002. Samplecollection for investigation of mars (SCIM): An early marssample return mission through the mars scout program. InProc. Lunar Planetary Sci. Conf. 33, abstract no. 1721.Huston: Lunar Planetary Institute.

158. Leshin, L. A., Clark, B. C., Forney, L. et al. 2003.Scienti�c Bene�t of a mars dust sample capture and earthreturn with SCIM. In Proc. Lunar Planetary Sci. Conf. 34,abstract no. 1288. Huston.

159. Mars exploration rovers.http://marsrovers.jpl.nasa.gov/overview/.

160. Eisen, H. cJ., Wen, L.C., Hickey, G. S., and Braun,D., 1998. Sojourner Mars Rover Thermal Performance.http://hdl.handle.net/2014/20229.

161. Where Are The Rovers Now?http://marsrovers.jpl.nasa.gov/mission/traverse_maps.html.

162. Paul1, H. L., and Diller, K. R. 2003. Comparison ofthermal insulation performance of �brous materials for theadvanced space suit. J. Biomech. Eng. 125: 639–647.

163. Fesmire, J. E., and Sass, J. P. 2008. Aerogelinsulation applications for liquid hydrogen launch vehicletanks. Cryogenics 48: 223–231.

164. Aspen Aerogels. http://www.aerogel.com/.

165. Braun, R. D., and Manning, R. M. 2007. Marsexploration entry, descent, and landing challenges.J. Spacecr. Rockets 44: 310–323.

166. Randall, J. P., Meador, M. A. B., and Jana, S. C.2011. Tailoring mechanical properties of aerogels foraerospace applications. ACS Appl. Mater. Interfaces 3:613–626.

167. Aerogel Particles/Cabot Aerogel.

168. Fesmire, J. E. 2006. Aerogel insulation systems forspace launch applications. Cryogenics 46: 111–117.

169. Fesmire, J. E., and Sass, J. P. 2008. Aerogelinsulation applications for liquid hydrogen launch vehicletanks. Cryogenics 48: 223–231.

170. Sumner, T. J. 2009. The STEP and GAUGE missions. SpaceSci. Rev. 148: 475–487.

171. Jones. S. M. 2006. Aerogel: Space explorationapplications. J. Sol-Gel Sci. Technol. 40: 351–357.

172. Visions and Voyages for Planetary Science 2013–2022.

173. Young, M. 2004. Investigation of several importantphenomen aassociated with the development of KnudsenCompressors. PhD dissertation. Los Angeles, CA: Universityof Southern California.

174. Han, Y.-L. 2006. Investigation of Knudsen Compressorsat low pressures. PhD dissertation. Los Angeles, CA:University of Southern California.

175. Basics of the Analysis Procedure.

http://solarsystem.nasa.gov/scitech/display.cfm?ST_ID=2402.

176. Pauzauskiea, P. J., Crowhursta, J. C., Worsley, M. A.et al. 2011. Synthesis and characterization of ananocrystalline diamond aerogel. PNAS 1010600108: 1–4.

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