light metals and their alloys ii : technology, microstructure and properties

Light Metals and their Alloys II

Technology,

Microstructure and Properties

Edited by Anna J. Dolata

Maciej Dyzia

Light Metals and their Alloys II

Technology, Microstructure and Properties

Special topic volume with invited peer reviewed papers only.

Edited by

Anna J. Dolata and Maciej Dyzia

Copyright 2012 Trans Tech Publications Ltd, Switzerland

All rights reserved. No part of the contents of this publication may be reproduced or transmitted in any form or by any means without the written permission of the publisher.

Trans Tech Publications Ltd Kreuzstrasse 10 CH-8635 Durnten-Zurich Switzerland http://www.ttp.net

Volumes 191 of Solid State Phenomena ISSN 1662-9787 (Pt. B of Diffusion and Defect Data - Solid State Data (ISSN 0377-6883))

Full text available online at http://www.scientific.net

Distributed worldwide by and in the Americas by

Trans Tech Publications Ltd Trans Tech Publications Inc. Kreuzstrasse 10 PO Box 699, May Street CH-8635 Durnten-Zurich Enfield, NH 03748 Switzerland USA

Phone: +1 (603) 632-7377 Fax: +41 (44) 922 10 33 Fax: +1 (603) 632-5611 e-mail: [email protected] e-mail: [email protected]

Introduction Faculty of Materials Engineering and Metallurgy was established in 1966 and currently is one of the 13 Faculties of Silesian University of Technology, located in Katowice. At present the faculty structure includes four departments: Metallurgy, Materials Technology, Materials Science and Management and Computer Science. The Faculty employs 38 professors and associate professors as well as 120 doctors (PhD). Scope of research activities includes materials engineering and metallurgy. The works carried out at the faculty are focused on research and development of advanced materials and their potential applications. Many scientific investigations are connected with problems of new technologies, formation the structure and properties of lightweight materials. This is the next collection of 30 articles presenting the results of research in scope of light metal alloys. That issue include three chapters: I – aluminium alloys, II – magnesium alloys and III – titanium alloys. Chapter I presents the subjects relating to the manufacturing of aluminum alloys, grain refinement and welding joints. This chapter presents also result of investigations concerning methods of obtaining and properties of aluminium matrix composites. Chapter II contain the papers presenting the results of researches carried out on conventional and new casting magnesium alloys. The first group of articles concern the effects of modification on the structure and properties of casting alloys. Following papers present results of researches on plastic deformation of Mg alloys. Subsequent articles cover topics related to the welding technologies. Last part of the chapter concern the magnesium matrix composites. Results of researches carried out on new generation of titanium alloys are presented in Chapter III. Papers included in this section concern the microstructure and properties Ti-Al base alloys. As well, possibilities of heat treatment and diffusion brazing of Ti alloys are discussed. This project is the second in the series of volume in the range of light metal alloys. The authors are planning to continue the series and publish every year. Editors.

Table of Contents

Introduction

Chapter 1: Aluminium and Aluminium Alloys

Numerical and Physical Modelling of Aluminium Refining Process Conducted in URO-200ReactorM. Saternus and T. Merder 3

Hydrodynamics of the Aluminium Barbotage Process Conducted in a Continuous ReactorM. Saternus 13

Influence of Overheating Degree on Material Reliability of A390.0 AlloyJ. Piątkowski 23

Mechamism of Grain Refinement in Al after COT DeformationK. Rodak and J. Pawlicki 29

Deformation-Induced Grain Refinement in AlMg5 AlloyK. Rodak, J. Pawlicki and M. Tkocz 37

CMT and MIG-Pulse Robotized Welding of Thin-Walled Elements Made of 6xxx and 2xxxSeries Aluminium AlloysJ. Adamiec, T. Pfeifer and J. Rykała 45

Fabrication of Ceramic-Metal Composites with Percolation of Phases Using GPIA. Boczkowska, P. Chabera, A.J. Dolata, M. Dyzia, R. Kozera and A. Oziębło 57

Producing of Composite Materials with Aluminium Alloy Matrix Containing SolidLubricantsA. Posmyk and J. Myalski 67

Machinability of Aluminium Matrix CompositesJ. Wieczorek, M. Dyzia and A.J. Dolata 75

Influence of Particles Type and Shape on the Corrosion Resistance of Aluminium HybridCompositesA.J. Dolata, M. Dyzia and W. Walke 81

Course of Solidification Process of AlMMC – Comparison of Computer Simulations andExperimental CastingR. Zagórski, A.J. Dolata and M. Dyzia 89

Chapter 2: Magnesium and Magnesium Alloys

Plasticity and Microstructure of Hot Deformed Magnesium Alloy AZ61D. Kuc, E. Hadasik and I. Bednarczyk 101

Effect of Modification on the Structure and Properties of QE22 and RZ5 Magnesium AlloysS. Roskosz, B. Dybowski and J. Paśko 109

Influence of Mould Cooling Rate on the Microstructure of AZ91 Magnesium Alloy CastingsS. Roskosz, B. Dybowski and R. Jarosz 115

Fractography and Structural Analysis of WE43 and Elektron 21 Magnesium Alloys withUnmodified and Modified Grain SizeS. Roskosz, B. Dybowski and J. Cwajna 123

Precipitate Processes in Mg-5Al Magnesium AlloyA. Kiełbus 131

Influence of Pouring Temperature on Castability and Microstructure of QE22 and RZ5Magnesium Casting AlloysB. Dybowski, R. Jarosz, A. Kiełbus and J. Cwajna 137

The Influence of Section Thickness on Microstructure of Elektron 21 and QE22 MagnesiumAlloysM. Stopyra, R. Jarosz and A. Kiełbus 145

b Light Metals and their Alloys II

The Influence of Tin on the Microstructure and Creep Properties of Mg-5Al-3Ca-0.7Sr-0.2Mn Magnesium AlloyT. Rzychoń and B. Chmiela 151

On the Oxidation Behaviour of WE43 and MSR-B Magnesium Alloys in CO2 AtmosphereR. Przeliorz and J. Piątkowski 159

Galvanic Corrosion Test of Magnesium Alloys after Plastic FormingJ. Przondziono, W. Walke and E. Hadasik 169

Creep Resistance of WE43 Magnesium Alloy JointsA. Kierzek and J. Adamiec 177

Impact of Heat Treatment on the Structure and Properties of the QE22 Alloy Welded JointsA. Kierzek and J. Adamiec 183

Microstructure of In Situ Mg Metal Matrix Composites Based on Silica NanoparticlesA. Olszówka-Myalska, S.A. McDonald, P.J. Withers, H. Myalska and G. Moskal 189

Microstructure of Mg-Ti-Al Composite Hot Pressed at Different TemperatureA. Olszówka-Myalska, R. Przeliorz, T. Rzychoń and M. Misiowiec 199

Chapter 3: Titanium and Titanium Alloys

The Chemical Composition and Microstructure of Ti-47Al-2W-0.5Si Alloy Melted inCeramic CruciblesW. Szkliniarz and A. Szkliniarz 211

Grain Refinement of Ti-48Al-2Cr-2Nb Alloy by Heat Treatment MethodA. Szkliniarz 221

Characteristics of Corrosion Resistance of Ti-C AlloysA. Szkliniarz and R. Michalik 235

Effect of a High-Temperature Hydrogen Treatment on a Microstructure and SurfaceFracture in Titanium Ti-6Al-4V AlloyM. Sozańska 243

Diffusion Brazing of Titanium via Copper LayerM. Różański and J. Adamiec 249

CHAPTER 1:

Aluminium and Aluminium Alloys

Numerical and physical modelling of aluminium refining process conducted in URO-200 reactor

Mariola Saternus1, a, Tomasz Merder1, b 1Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

a [email protected], b [email protected]

Keywords: aluminium refining, physical modelling, numerical modelling.

Abstract. At present both primary and secondary aluminium needs to be refined before further

treatment. This can be done by barbotage process, so blowing small bubbles of inert gas into liquid

metal. This way harmful impurities especially hydrogen can be removed. Barbotage is very

complex taking into consideration hydrodynamics of this process. Therefore modelling research is

carried out to get to know the phenomena that take place during the process better. Two different

modelling research can be applied: physical and numerical. Physical modelling gives possibility to

determine the level of gas dispersion in the liquid metal. Whereas, numerical modelling shows the

velocity field distribution, turbulent intensity and volume fraction of gas.

The paper presents results of physical and numerical modelling of the refining process taking place

in the bath reactor URO-200. Physical modelling was carried out for three different flow rate of

refining gas: 5, 10 and 15 dm3/min and three different rotary impeller speeds: 0, 300, 500 rpm

Commercial program in Computational Fluid Dynamics was used for numerical calculation. Model

VOF (Volume of Fluid) was applied for modelling the multiphase flow.

Obtained results were compared in order to verify the numerical settings and correctness of the

choice.

Introduction

Today, both primary and secondary aluminium contains many impurities such as hydrogen or

nonmetallic and metallic inclusions. Hydrogen content in aluminium and its alloys is in the range

between 0.05 to 0.6 cm3/100g Al [1]. To reduce hydrogen concentration to the level 0.06 – 0.07

cm3/100g Al refining process is applied. Additionally, parts of nonmetallic and metallic inclusions

can be simultaneously eliminated by means of flotation. Therefore, aluminium refining process has

become one of the integral technological stages in obtaining aluminium. The most commonly used

method is barbotage that means blowing the inert gases, especially argon into the liquid aluminium.

There are different kinds of refining reactors: batch and continuous. Small gas bubbles are

generated by ceramic porous plugs, different kind of nozzles and rotary impellers. All over the

world there are many technological solutions of such reactors, for example: ACD, AFD, Alcoa 622,

ASV, DMC, DUFI, JetCleaner, GBF, GIFS, Hycast, LARS, MINT, RDU, Rotoxal, Shizunami

[2,3]. In Poland one of the most popular reactors is the URO-200 reactor designed by IMN-OML in

Skawina. This reactor works in many polish foundries. The problem connected with this type of

reactor is obtaining the uniform dispersion of gas bubbles in the whole volume of liquid. The

influence on this has the following processing parameters: flow rate of refining gas and mainly the

rotary impeller speed. The choice of these parameters allows to optimize the industrial aluminium

refining process.

Main information about modelling

Generally, the barbotage process is characterized by high dynamics of course because of the quick

mass transfer between phases. So, it is essential to understand the phenomena occurring during the

process and determine the hydrodynamic conditions in which the process takes place. They

influence directly the value of mass exchange area and the mass transfer coefficient. Obtaining the

© (2012) Trans Tech Publications, Switzerlanddoi:10.4028/www.scientific.net/SSP.191.3

appropriate size of gas bubbles and their dispersion in the liquid metal ensures good efficiency of

the process. The process of gas bubbles creation and their movement is very complex, so its

analytical description present fundamental difficulties. As a solution in this case, modellling

research is applied. In metallurgy there are many methods of modelling the liquid flow (see Fig. 1).

Fig. 1. Modelling of the liquid flows applied in metallurgy [4]

The hydrodynamic conditions can be determined by the physical modelling. However, the flow of

mass and gas is not fully shown by this modelling, these kinds of research are very often and

willingly [5-9] used due to difficulties in conducting experimental test in real conditions (sometimes

this is even impossible). Additionally, modelling research is not as expensive as the one carried out

in industrial conditions. This method gives possibility to obtain information about phenomena

occurring in the liquid metal during the blowing process of gas bubbles. Water is used as a

modelling agent of aluminium. It is applied because its accessibility, low costs, and especially the

fact that some physical features of water in room temperature are similar to features of aluminium

in temperature 700 0C (e.g. dynamic viscosity).

If the results obtained from this kind of research are to be representative and can be transferred into

the real conditions, the physical model has to be built according to the strict rules coming from the

theory of similarity [10-13]. This similarity concerns the characteristic features of the real object

that have important influence on the phenomena occurring in the examined process. Taking into

account the construction of the examined metallurgical units it is essential to fulfill the following

conditions:

geometric similarity of the model and a real object,

hydrodynamic similarity for the liquid flow in the model and the object which especially

concerns: kinetic similarity, dynamic similarity, heat similarity.

Fulfillment of the similarity rules according to the theory of dimensional analysis, can be done

basing on the equality rule of the appropriate criterial numbers in the model and the examined

object. The results obtained from the experimental test on the physical model, after verification, can

be transferred to the real conditions. Taking into account both - the construction of refining reactors

used for the barbotage process and the hydrodynamics of the metal flow, the adequate criterial

4 Light Metals and their Alloys II

numbers describing the process are: Euler`s number, Reynold`s number, Froude`s number and

Weber`s number. Table 1 presents the values of Reynold`s, Weber`s and Froude`s numbers for

water (in 293 K) and aluminium (in 973K) for the batch refining reactor URO-200.

Table 1. Values of he criterial numbers values for water (in 293 K) and aluminium (in 973K)

Criterial number Reynold`s number Weber`s number Froude`s number

Value water 27802.0 84.24 0.0029

aluminium 67392.0 21.41 0.0029

The other kind of modelling is numerical simulation. Numerical and physical modelling

complements each other, and as a consequence the analyzed process can be understood better. It

additionally, allows to obtain useful information such as: determining the level of distribution or

participation of gaseous phase. The conducted numerical calculations have to correspond the

conditions in which the experimental tests on water model of refining reactors were carried out. So,

numerical analysis requires choosing the proper model describing the physical phenomena

occurring in the process. The introduction of gas into the liquids is a reason why the problem is

considered as a multiphase (diphase) flow. It is possible to take into consideration the real object or

the examined model (argon-aluminium, argon-water). In case of multiphase flows there are many

ways to describe the process mathematically. One solution of the diphase flow problems is

described by Langrange. Then, the movement equations are solved directly for every particle. This

method is however very time-consuming. The more effective, taking into account the calculating

time, is description of the diphase flow by means of Euler`s method. This method on the other hand

requires the application of simplified assumptions and the appropriate formulation of the boundary

conditions [14]. So, the choice of model is always a compromise between the accuracy of solutions,

and the required calculating time [15].

Complexity of mathematical model and requirements of calculation correctness enforce the

application of the effective tool for numerical solution of the system of partial differential

equations. In this case the most appropriate seems to be commercial code AnsysFluent. It is

equipped with the numerical procedure set, needed to solve the system of modelling equations after

choosing initial and boundary conditions. The numerical solution is based on the control volume

method. In this method the partition of the calculating space of the object is made by means of the

net with the particular number of cells (control volumes) [16]. In multiphase flows (in this case

diphase flow) two methods of the solution are applied: Euler-Lagrange`s method and Euler-Euler`s

method (see Fig. 2).

Fig. 2. Division of methods used for solving the diphase flow in the Fluent program

Measuring apparatus and experimental procedure

Research concerning physical modelling was carried out for the batch refining reactor URO-200.

This reactor was designed and constructed in the Institute of Non-Ferrous Metals – Light Metals

Division in Skawina. The model for this reactor was built at 1:1 scale. Conditions that can be

Anna J. Dolata and Maciej Dyzia 5

observed in a real reactor correspond to the test stand used for modelling research. The thermal

expansion of gas was not considered. It was made an assumption that there are isothermal

conditions that means during the gas mixing there is no considerable differences of temperatures.

Fig. 3 shows the test stand used for modelling research. Argon was used as a refining gas and water

as a modelling agent. During tests the flow rate of the refining gas and rotary impeller speed were

changing. Every case was registered by the digital camera.

Fig. 3. a) The test stand used for modelling research b) with the description

For numerical modelling the studied problem was simplified to the 2D object in order to choose the

model of multiphase flow and check correctness of accepted simplifications and boundary

conditions. Target mesh consists of 13180 controlled volumes. Fig. 4 presents the scheme of mesh

and applied boundary conditions. Taking into account the symmetry of the object, numerical

calculations were carried out for the half of the object.

Fig. 4. Scheme of the object taken to numerical calculations: a) mesh, b) boundary conditions

Numerical simulations were carried out by means of Volume of Fluid model (1 phase - water, 2

phase – argon). The flow in boundary layer was modelled using the Standard Wall Function. This

method let to decrease the calculation inputs considerably. Analytical solution was used in the

boundary area to describe the velocity fields, so in this area fewer numbers of nodes could be used.


In calculations the algorithm PISO, which is recommended by the producer, was applied. For

digitization of the pressure the scheme PRESTO! was used. Numerical procedures of the second

order upwind were applied. Calculations were made in unsteady conditions using time step size ∆t=

0.001s. In numerical calculations three cases of different flow rate of refining gas without rotation

were analyzed. Table 2 shows data concerning modelling variants of the flow of refining gas and

the appropriate gas bubble diameter.

Table 2. Data concerning modelling variants of the flow of refining gas and the gas bubble diameter

Parameter of the process

Flow rate of refining gas, [l/min] 5 10 15

Mass flow of refining gas, [kg/s] 1.4e-4 2.7e-4 4.1e-4

Gas bubble diameter, [m] 0.003 0.005 0.006

Physical modelling – results of the research

The research was carried out in the Department of Metallurgy at the Silesian University of

Technology. Influence of the flow rate of refining gas and rotary impeller speed on the level of gas

dispersion in water was examined. The flow rate of refining gas was changing in the range from 5

to 15 dm3/min every 5 dm

3/min. Research was conducted without rotation and with rotation for two

different rotary impeller speeds: 300 rpm and 500 rpm. Fig. 5 presents registered results for three

different flow rate of refining gas without rotation. It can be seen that smaller or bigger single gas

bubbles raise up to the top of the reactor. Dispersion occurs only in the area in which the gas

bubbles are generated – this place is marked with black rectangle on the pictures. There is no

dispersion in the whole volume of liquid, so the case of minimal dispersion is observed.

a) b) c)

Fig. 5. Results of blowing argon into the liquid for different flow rate of gas: a) 5 dm3/min,

b) 10 dm3/min, c) 15 dm

3/min without rotation

Fig. 6 presents registered results for three different flow rates of refining gas with rotary impeller

speed equaled 300 rpm. Single gas bubbles raise up to the top of the reactor. Gas bubbles are well

mixed with the liquid – the area of dispersion is marked with the black rectangle. Sometimes only

near the side walls of the reactor there is a lack of dispersion. Swirls on the liquid surface make the

bubbles in the upper part of the reactor mix with liquid. Generally the case of uniform dispersion,

which is the most desirable, is observed.


Fig. 7 presents registered results for three different flow rates of refining gas with rotary impeller

speed equal 500 rpm. Single gas bubbles raise to the top of the reactor and somewhere they create

chains. Gas bubbles are well mixed with the liquid. In some parts of the reactor – marked by the

arrow - there is a chain flow of refining gas, so there are good conditions for the existence of swirls.

It means that there is some danger to introduce again hydrogen from the surface to the metal.

a) b) c)

Fig. 6. Results of blowing argon into the liquid for different flow rates of gas: a) 5 dm3/min,

b) 10 dm3/min, c) 15 dm

3/min with the rotary impeller speed equaled 300 rpm

a) b) c)

Fig. 7. Results of blowing argon into the liquid for different flow rates of gas: a) 5 dm3/min,

b) 10 dm3/min, c) 15 dm

3/min with the rotary impeller speed equaled 500 rpm

Numerical modelling – results of research

Calculations were done for different flow rate of blown argon (5, 10 and 15 dm3/min). As a result,

expected velocity field distribution, turbulent intensity and volume fraction of argon were obtained.

Results are presented for three periods (after 2, 10 and 30 s). After 30 s the steady conditions of the

process are reached. Fig. 8 presents forecasted velocity field distribution of water for different mass

flow of injected argon after different time. Additionally, the steam lines were put in order to picture

the movement of liquid (water) better. It can be seen that the movement of liquid went on from

rotary impeller to the side walls of the crucible and then came back to the rotary impeller (their

circuit). Results of velocity field distribution show that the flow of the mixture in the object is

dominated by the injected gas. Basing on the analysis of mixture flow it has been found that the


injected gas is responsible for creating swirls. This is confirmed by the turbulent intensity of water

for the examined flow rate of argon (see Fig. 9). The biggest kinetic energy is observed near the

metal surface (sometimes in real condition when there is no rotation small gas geysers are observed,

especially when the flow rate of gas is high and the gas bubbles are big one – see Fig. 10). These

results are complementary to the physical modelling research.

Fig. 8. Results of calculations – velocity field distribution [m/s] with the steam lines for different

argon mass flow: a) 1.4e-4 kg/s, b) 2.7e-4 kg/s, c) 4.1e-4 kg/s

Verification of the model

Fig. 11 presents the volume fraction of argon for different flow rates of refining gas: 5, 10 and 15

dm3/min. To compare obtained results from numerical modelling with results obtained from

physical modelling they were juxtaposed. It can be noticed that there is coincidence between results

coming from both types of modelling. The same trajectory of gas bubbles movement can be


observed. However, the number of gas bubbles in the presented schemes does not agree with the

number of gas bubbles obtained in the industrial conditions (3D object). This can be explained by

simplifying the real object to two-dimensional object.

Fig. 9. Results of calculations – turbulent kinetic energy [m

2/s

2] for different argon mass flow:

a) 1.4e-4 kg/s, b) 2.7e-4 kg/s, c) 4.1e-4 kg/s

Summary

The effectiveness of hydrogen removal process from liquid aluminium depends not only on the flow

rate of refining gas but also the rotary impeller speed. The determination of the optimal values of

such parameters allows to reduce the costs of the process. So the most desirable level of gas

dispersion in the liquid metal is obtained when the gas bubble mixing with the liquid is observed in

the whole volume of the reactor. However if the flow rate of rotary impeller speed is too high then

the chain flow and swirls in the liquid metal and on the metal surface can be seen. The most optimal

seems to be the flow rate of refining gas between 10 and 15 dm3/min and the rotary impeller speed

300 rpm. If the rotary impeller speed is higher (500 rpm) then swirls can be created. These swirls

can be the reason why hydrogen is introduced to the metal again.


Numerical modelling allows to obtain velocity field distribution, turbulent intensity and volume

fraction of argon in a liquid. Convergence of results obtained from two different methods confirms

the choice of VOF model for numerical simulation and also that the assumptions were made

properly. Thus results obtained from numerical modelling can be used in estimating the phenomena

that take place during the aluminium refining process, especially barbotage. The optimal seems to

a) b) c)

Fig. 10. a) Scheme of formation of metal geysers on the metal surface during the barbotage process

without rotation; results coming from physical modelling: b) no rotation, argon flow rate: 15

dm3/min, the view of the surface c) no rotation, argon flow rate: 20 dm

3/min, big gas bubbles,

bigger geyser on the surface

Fig. 11. Volume fraction of argon for different argon mass flow: a) 5 dm

3/min (1.4e-4 kg/s), b) 10

dm3/min (2.7e-4 kg/s), c) 15 dm

3/min (4.1e-4 kg/s) obtained from numerical calculation and

comparison with results coming from physical modelling


be joining physical and numerical modelling. They complement and verify each other, so the

application of the obtained results in industrial conditions is easier. The next step of research should

focus on numerical simulation for three-dimensional case. However this is not easy to obtain

because of the long calculation time and applying moving net for the calculation. On the other hand

the preliminary results of the research are very promising. They may complement the results

obtained for the physical model.

Acknowledgements to the State Committee for Scientific Research (KBN-MNiSW) for financial

support (project Nr N N508 443236).

References

[1] X.-G. Chen, F.-J. Klinkenberg, S. Engler, Optimization of the impeller degassing process

trough continuous hydrogen measurements, Light metals TMS (1995) 1215-1222.

[2] M. Saternus, J. Botor, Aluminium refining process – methods and mathematical models,

ALUMINIUM 81 (2005) 209-216.

[3] M. Saternus, J. Botor, Refining of aluminium and its alloys by inert gases, Ores and

Nonferrous metals 48 (2003) 154-160.

[4] H.-J. Odenthal, R. Bölling, H. Pfeifer, Numerical and physical simulation of tundish fluid flow

phenomena, Steel Research 74 (2003) 44-55.

[5] J.W. Evans, A. Field, N. Mittal, Measurements of bubble dispersion and other bubble

parameters in a gas fluxing unit at Alcoa using a capacitance probe, Light metals TMS (2003)

909-913.

[6] S.T. Johansen, S. Graadahl, P. Tetlie, B. Rasch, E. Myrbostad, Can rotor-based refining units

be developed and optimized based on water model experiment, Light Metals TMS (1998)

805-810.

[7] E. Waz, J. Carre, P. Le Brun, A. Jardy, C. Xuereb, D. Ablitzer, Physical modelling of the

aluminium degassing process: experimental and mathematical approach, Light metals TMS

(2003) 901-907.

[8] K.A. Carpenter, M.J. Hanagan, Efficiency modeling of rotary degasser head configuration and

gas introduction methods, Part 1 – water tank test, Light metals TMS (2001) 1017-1020.

[9] J.L. Song, F. Chiti, W. Bujalski, A.W. Nienow, M.R. Jolly, Study of molten aluminium

cleaning process using physical modelling and CFD, Light Metals TMS (2004) 743-748.

[10] L. Müller, Dimensional Analysis Applying in Research of Models, PWN, Warsaw, 1983.

[11] K. Michalek, Physical and Numerical Modelling Using for Optimization of Metallurgical

Processes, VSB – Technical University of Ostrava, Ostrava, 2001.

[12] K. Michalek, K. Gryc, M. Tkadlečková, D. Bocek, Model study of tundish steel intermixing

and operational verification, Archives of metallurgy and materials 57 (2012) 291-296.

[13] K. Michalek, M. Tkadlečková, K. Gryc, P. Klus, Z. Hudziczek, V. Sikora, P. Střasák,

Optimization of argon blowing conditions for the steel homogenization in ladle by numerical

modelling, METAL (2011) 143-149.

[14] R. Grybos, The Basic of Fluid Dynamics – part II, PWN, Warsaw, 1998.

[15] T. Merder, J. Pieprzyca, Numerical modeling of the influence subflux controller of turbulence

on steel flow in the tundish, Metalurgija 4 (2011) 223-226.

[16] FLUENT, User's Guide, version 6.0, Fluent Inc., 2003.


Hydrodynamics of the aluminium barbotage process conducted in a continuous reactor

Mariola Saternus1, a 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

a [email protected]

Keywords: aluminium refining, barbotage process, physical modelling.

Abstract. Today aluminium obtained from ores (primary) and from scrap (secondary) need to be

refined. During this process harmful impurities such as hydrogen, sodium, lithium, oxides, borides

or carbides can be removed. There are many different ways of aluminium refining process. The

most popular seem to be barbotage that means blowing through aluminium many tiny gas bubbles

of refining gas. Reactors applying this methods have been working all over the word. They are of

different types: bath and continuous, using ceramic porous plugs, special kinds of nozzles or rotary

impeller for generating small gas bubbles. At present reactors for continuous refining have become

the most popular. In Poland typical representative of such reactors is URC-7000 reactor.

The phenomena occurring during this process are rather complicated. Therefore to know them

better the modelling research is applied, especially physical modelling. The paper presents the

results of such a research. The tests were carried out in the test stand for modelling the barbotage

process in the URC-7000 reactor. The different modelling agents were tested (water, glycerin and

mixture of water and glycerine). The density and viscosity of water and glycerin mixture were

determined. Modelling tests were conducted for four different flow rates of refining gas: 6, 10, 15

and 20 dm3/min. Results were registered by digital camera. Pictures for different modelling agents

were juxtaposed and discussed.

Introduction

Today aluminium can be obtained from ores via Bayer`s method and electrolysis process of Al2O3

or from scrap via recycling. Aluminium obtained in such ways aluminium contains many impurities

such as hydrogen, metallic inclusions (Na, Li) and nonmetallic inclusions (oxides, borides, nitrides,

carbides). These impurities in both primary and secondary aluminium make the properties of metal

worse. This concers especially the creation of porosity which negatively influences the mechanical

properties (see Fig. 1) [1,2]. Therefore, refining process has become an important stage of obtaining

aluminium. There are many technological solution for removing impurities from aluminium but

recently the most popular has become barbotage process that means blowing the liquid metal by

small gas bubbles of refining gas [3,4]. This process can be conducted in bath reactors but more and

more popular are continuous ones. The examples of continuous reactor are following: ACD, AFD,

Alcoa 469, Alcoa 622, Alpur, DMC, DUFI, FIF-50, FILD, GBF, GIFS, HYCAST, I-60 SIR,

Jetclenaer, LARS, MINT, RDU. The refining gas can be introduced by special nozzles, ceramic

porous plugs and rotary impellers. In Poland typical representative of this kind of reactor is URC-

7000 (see Fig. 2). It was designed in Skawina in the Institute of Nonferrous Metals – Light Metal

Division. Flow rate of refining gas for this reactor is in the range from 100 to 320 dm3/min. Table 1

presents main parameters of continuous reactor working all over the world.

Physical modelling

Modelling research (physical and numerical) is commonly used for analyzing and learning about

the phenomena occurring in reactors which are applied in metallurgy of steel and nonferrous metals

[7-11]. Difficulties in conducting experimental test in real conditions are the reason why physical

modelling is applied much more often.


Fig. 1. Change of mechanical properties with the increasing content of porosity in aluminium alloy

CP601: a) UTS, b) elongation [1,2]

Fig. 2.View of URC-7000 reactor [5]

Table 1. Main parameters of continuous reactors used for aluminium barbotage process [6]

Reactor Flow rate of

metal

[kg/h]

Final hydrogen

concentration

[cm3/100 g Al]

Reactor Flow rate of

metal

[kg/h]

Final hydrogen

concentration

[cm3/100 g Al]

ACD 15000 – 19500 0.12 – 0.10 Alcoa 622 9000 – 13600 0.22 – 0.15

AFD 19800 0.14 – 0.10 Alpur 35000 – 60000 0.45 – 0.10

DMC 19200 – 40800 0.11 – 0.06 DUFI 2500 – 20000 0.19 – 0.08

FILD 3800 0.13 – 0.04 GBF 12000 – 42000 0.12 – 0.05

GIFS 18000 – 39000 0.16 – 0.07 HYCAST 20000 – 60000 0.09

I-60 SIR 10000 – 65000 0.13 – 0.10 LARS 8100 – 22500 0.12 – 0.09

MINT 5000 – 15000 0.25 – 0.05 RDU 30000 0.30 – 0.05

SNIF 27000 – 36000 0.25 – 0.07 URC-7000 5000 0.10

Additionally, it allows to gain information about hydrodynamic phenomena occurring during the

process. What is more, modelling research is not as expensive as the one carried out in industrial

conditions. If the results obtained from this kind of research are to be representative and can be

transferred into the real conditions, the physical models have to be built according to the fix rules

coming from the theory of similarity [12-15]. This similarity concerns the characteristic features of

the real object that have important influence on the phenomena occurring in the examined process.

So, the following conditions have to be fulfilled:

• similarity of the model and a real object,

• hydrodynamic similarity for the liquid flow in the model and the object (kinetic similarity,

dynamic similarity, heat similarity).


In physical model this problem can be solved with the help of the appropriate characteristic criterial

numbers. For aluminium barbotage process the most important criterial numbers are: Euler`s,

Reynold`s, Froude`s and Weber`s number.

Euler`s number is considered as a ratio of pressure differences in defined two points of model to the

dynamic pressure. The value of Euler`s number is most often searching for, so it is seen as a

dependent variable, which can be presented as a relationship of other criterial number:

Eu = f (Re, Fr). (1)

Euler`s number has great significance when there is a case of flow under pressure, in other cases

(open channels or reactors) it can be neglected.

Reynold`s number is treated as a ratio of the dynamic forces to the friction force occurring in the

flow of liquid:

Re = ρ·u·L·η-1

. (2)

where: ρ – density, u – flow velocity, L – characteristic dimension, η – coefficient of dynamic

viscosity.

When the laminar flow is observed the values of Reynold`s number are small, however when these

values are big the turbulent flow is quoted. Very often transfer from laminar flow to turbulent flow

is violent, so then the limiting value of Reynold`s number is determined as a critical value of Re. In

that range of flows the value of Reynold`s number changes insignificantly (if the character of flow

is not changing). This range is know as a selfmodelling region that means region in which the

studied phenomenon is practically independent of the similarity numbers, so in that case there is no

necessity to obtain the equality of criterial numbers.

Froud`s number is considered as a ratio of the dynamic forces to the terrestrial gravity forces:

Fr = u2·g

-1·L

-1. (3)

where: g – gravitational acceleration.

The Froude`s criterion has to be taken into account when modelling process takes place in the

reactor in which the gravitational forces are important. Usually the same gravitational acceleration

influences on the model and analyzed object. So, the Froude`s similarity can be presented in the

following form:

Fr = u2·L

-1 = u`

2·L`

-1 = Fr`. (4)

which means that in the scaled-down model appropriately smaller velocity of liquid flow should be

applied and scales are fulfilled according the proportion:

Su = SL0.5

. (5)

The regulation of flow velocity can be done taking into account the change of flow rate of

modelling agent. So, for the scale of flow rate the following relation can be written:

SQ = Su · SL2 = SL

5/2. (6)

Weber`s number is a ratio of the force of inertia to the force of surface tension:

We = u2·L·ρ·σ

-1. (7)

where: σ – surface tension.

In physical modelling as a typical modeling agent water is used because of its accessibility, low

costs, and especially the fact that some physical features of water in room temperature are similar to

features of aluminium in temperature 700 0C (e.g. dynamic viscosity). Sometimes also glycerine is

used. Table 2 shows the characteristic parameters of aluminium, glycerin and water needed for

calculation criterial numbers and values of these numbers (Reynold`s, Weber`s and Froude`s

number) for the continuous refining reactor URC-7000.


Table 2. Basic parameters of aluminium, glycerin and water and values of the criterial numbers

Characteristic parameters/number aluminium glycerine water

Temperature [K] 973 293 293

Dynamic viscosity [Pa·s] 0.00100 0.93400 0.00101

Surface tension [N·m-1] 0.680 0.063 0.072

Density [kg·m-3] 2400 1260 1000

Reynold`s number 4620.0 2.6 1905.9

Weber`s number 0.118 0.669 0.467

Froude`s number 0.00028 0.00028 0.00028

Measuring apparatus and experimental procedure

The research was carried out in the Department of Metallurgy at the Silesian University of

Technology. Research concerning physical modelling was conducted for the continuous refining

reactor URC-7000. The model for this reactor was built at 1:3 scale. Conditions that can be

observed in a real reactor match the test stand used for modelling research. Fig. 3 shows the scheme

of test stand used for modelling research and Fig. 4a the real view of this test stand.

Argon was used as a refining gas. For modelling the thermal expansion of gas was not considered.

It was made an assumption that there are isothermal conditions that means there is no considerable

differences of temperatures during the gas mixing. Also the interfacial tension for creation of

bubbles was neglected because forces of surface tension are considered in the criterial numbers.

Water, glycerin and mixture of water and glycerin in different percentage friction (20, 40, 60, 80%)

were used as a modelling agent. The density and viscosity of water and glycerin mixture were

measured by means of areometers and Höppler`s viscometer (see Fig. 4b and c). During the tests,

the flow rate of refining gas was changing from 6 to 20 dm3/min. Every case was registered by the

digital camera.

Fig. 3. Scheme of test stand used in physical modelling

Fig. 4. a) Real view of test stand used in physical modelling; b) Areometers and c) Höppler`s

viscometer used for measuring density and viscosity


Results of the research

Table 3 presents results of density and viscosity measurement for different mixture of glycerin and

water. Basing on this, Reynold`s number was calculated (see also Table 3). Influence of the flow

rate of refining gas on the level of gas dispersion in water, glycerin and their mixtures were

examined. Fig. 5 presents registered results for different modelling agents when the flow rate of

refining gas equals 5 dm3/min whereas Fig. 6 to 8 show results for the flow rate 10, 15 and 20

dm3/min respectively.

Table 3. Values of density, viscosity and Reynold`s number of water and glycerin mixtures

Parameter/number glycerin 80% glycerin 60% glycerin 40% glycerin 20%

Density [kg/m3] 1207 1159 1114 1056

Viscosity [Pa·s] 0.04682 0.00823 0.00309 0.00145

Reynold`s number 49.6 271.1 694.0 1401.9

Summary

It can be seen that if in water model the flow rate of refining gas is smaller or equals 6 dm3/min

dispersion of gas bubbles in the liquid occurs only in the area in which gas bubbles are generated.

There is no dispersion in the whole volume of liquid, especially in the middle of the reactor between

ceramic porous plugs and in the bottom below the plugs. The same can be stated for glycerin and

mixture of glycerin and water. For glycerin and mixtures of glycerin and water smaller gas bubbles

are observed, however the results are similar. In pure glycerin (see Fig. 5b) gas bubbles slowly rise

to the surface.

The most optimal seems to be the flow rate of refining gas equal 10 dm3/min. For this case the gas

dispersion in the whole reactor is uniform. Only in the bottom near the plugs there is no dispersion.

For glycerin and mixture of water and glycerin the gas dispersion is observed in the whole reactor,

but the mechanism is different. At first the big gas bubbles raise to the surface as a chain but then

small gas bubbles spread to the whole volume of liquid. For the flow rate of refining gas 15

dm3/min the situation is almost the same, however here undesirable swirls on the surface can be

seen – especially in glycerin. The flow rate of refining gas equals 20 dm3/min is too high. Gas

bubbles are well dispersed in the whole volume, but on the surface waving can be seen, that means

swirls are created and as a consequence the hydrogen can be picked again by the metal. To sum up,

the mechanism of gas bubble spreading in the whole dispersion in the case of glycerin and water are

different. However, the obtained results are almost the same. The most optimal flow rate of refining

gas for the model seems to be between 10 and 15 dm3/min.


a) b)

c) d)

e) f)

Fig. 5. Results of blowing argon into the liquid for the flow rate of gas equal 6 dm3/min with the

following media as a modelling agent: a) water, b) glycerin, mixture of water and glycerin: c) 80%, d)

60%, e) 40%, f) 20%

So, taking into account the operational conditions of the reactor the flow rate of refining gas should

be between 150 to 250 dm3/min. In that case the process of uniform dispersion of refining gas in the

liquid aluminum can be obtained and then the level of hydrogen concentration will be at the

appropriate level (lower than 0.1 cm3/100 g Al). The higher flow rate of refining gas is not

appropriate from the economic point of view. Additionally, there is danger that when the chain flow

is observed hydrogen can be picked up from the atmosphere again to the metal and then the level of

hydrogen concentration will be higher than 0.1 cm3/100 g Al. This case is not desirable.


a) b)

c) d)

e) f)


following media as a modelling agent: a) water, b) glycerin, mixture of water and glycerin: c) 80%,

d) 60%, e) 40%, f) 20%


a) b)

c) d)

e) f)



d) 60%, e) 40%, f) 20%


a) b)

c) d)

e) f)

Fig. 8. Results of blowing argon into the liquid for flow rate of gas equal 20 dm3/min with the


d) 60%, e) 40%, f) 20%

Acknowledgements to the State Committee for Scientific Research (KBN-MNiSW) for financial

support (project Nr N N508 443236).


References

[1] A.M. Samuel, F.H. Samuel, Review various aspects involved in the production of low-

hydrogen aluminium castings, Journal of Materials Science 27 (1992) 6533-6563.

[2] J.A. Eady, D.M. Smith, Effect of porosity on the tensile properties of aluminium castings,

Mater. Forum 9 (1986) 217-223.

[3] M. Saternus, J. Botor, Aluminium refining process – methods and mathematical models,

ALUMINIUM 81 (2005) 209-216.

[4] Y. Liu, T. Zhang, M. Sano, Q. Wang, X. Ren, J.He, Mechanical stirring for highly efficient gas

injection refining, Trans. of Nonferrous Metals Society of China 21 (2011) 1896-1904.

[5] Technical Documentation of URC-7000 Reactor, Nicromet – Oświęcim, Skawina, 2003.

[6] M. Saternus, Refining Process of Aluminium and its Alloys by Means of Argon Blowing,

Silesian University of Technology ed., Gliwice, 2011.

[7] S.T. Johansen, S. Graadahl, P. Tetlie, B. Rasch, E. Myrbostad, Can rotor-based refining units be

developed and optimized based on water model experiment, Light Metals TMS (1998) 805-

810.

[8] E. Waz, J. Carre, P. Le Brun, A. Jardy, C. Xuereb, D. Ablitzer, Physcial modelling of the

aluminium degassing process: experimental and mathematical approach, Light metals TMS

(2003) 901-907.

[9] K.A. Carpenter, M.J. Hanagan, Efficiency modeling of rotary degasser head configuration and

gas introduction methods, Part 1 – water tank test, Light metals TMS (2001), 1017-1020.

[10] M. Warzecha, T. Merder, H. Pfeifer, J. Pieprzyca, Investigation of flow characteristics in a

six-strand CC tundish combining plant measurements – physical and mathematical modelling,

Steel Research International 81 (2010) 987-993.

[11] T. Merder, J. Pieprzyca, Numerical modeling of the influence subflux controller of turbulence

on steel flow in the tundish, Metalurgija 4 (2011) 223-226.

[12] L. Müller, Dimensional Analysis Applying in Research of Models, PWN, Warsaw, 1983.

[13] K. Michalek, J. Morávka, K. Gryc, Mathematical indentification of homogenisation processes in

argon stirred ladle, Metalurgija 48 (2009) 219-222.

[14] K. Michalek, K. Gryc, J. Moravka, Physical modelling of bath homogenization in argon stirred

ladle, Metalurgija 48 (2009) 215-218.

[15] K. Michalek, Z. Hudziczek, K. Gryc, M., Tkadleckova, Study of homogenization and transfer

processes in the casting ladle using physical modelling, METAL (2010) 42-46.


Influence of overheating degree on material reliability of A390.0 alloy

Jarosław Piątkowski1,a 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

a [email protected]

Keywords: overheating degree, hypereutectic alloy, statistics, Weibull distribution.

Abstract. The object of the studies was A390.0 alloy (AlSi17Cu5Mg), similar to A3XX.X series,

gravity cast into sand and metal mould. This alloy is mainly used for cast pistons operating in I.C.

engines, for cylinder blocks and housings of compressors, and for pumps and brakes. The A390.0

alloy was poured at temperatures 880 and 980oC, holding the melt for 30 minutes and cast from the

temperature of 780oC. The assessment of the impact of the degree of overheating was to analyses

the tensile strength. Studies were carried out on a normal-running fatigue testing machine, which

was the mechanically driven resonant pulsator. For the needs of quantitative reliability evaluation

and the time-to-failure evaluation, the procedures used in survival analysis, adapted to the analysis

of failure-free operation with two-parametric Weibull distributions, were applied. Having

determined the boundary value „σ0” for Weibull distribution, the value of „m” modulus was

computed along with other coefficients of material reliability, proposed formerly by the authors.

Basing on the obtained results, a model of Weibull distribution function was developed for the

tensile strength with respective graphic interpretation.

Introduction

Studies on the structure of liquid metals and alloys [1-5] do not allow yet a definitive and

unambiguous assessment of processes going on in these materials, but our knowledge of these

phenomena can have a significant effect on the determination of molten alloy predisposition to the

occurrence of certain type of crystalline structures after solidification and consequently to

improvement of casting properties.

Known from literature [6], the cluster theory used in the theoretical description of liquid metal as

well as other hypotheses (statistical, condensation, network, geometric, thermodynamic) [7] do not

fully explain numerous phenomena that take place in liquid metal. Therefore, it seems advisable to

undertake research on the effect of time and temperature of heat treatment on the tensile strength of

alloy. The results of these studies will make basis for further determination of alloy performance

reliability measured by Weibull modulus „m” [8-9].

Scope and aim of investigations

The main aim of the investigations was determined of Weibull modulus for A390.0 alloy. The

results of static tensile test were basis for the determination of a boundary value of „σ0” used in

Weibull distribution and of the value of modulus „m” generally considered a measure of the

performance reliability of the investigated alloy.

The scope of investigations included:

− determination of tensile strength (32 samples were cast for each temperature value),

− determination of the main estimators and variability indeces,

− calculation of Weibull modulus,

− development of graphical relationships for the survival probability „PS” in function of the tensile

stress „σ”, considered a measure of the performance reliability of examined alloy.


Test materials and methods

Selected for the tests A390.0 cast alloy of the chemical composition shown in Table 1.

Table 1. Chemical analysis of the A390.0 alloy.

Alloy Chemical composition, [wt.%]

Si Cu Fe Mn Mg Ni Al

A390.0 16,63 4,87 0,44 0,03 0,94 0,02 rest

The stand for melting and casting of tensile specimens to examine the silumin properties is shown

in Figure 1.

Fig. 1. Test stand for melting and casting of A 390.0 alloy.

Schematic representation of the concept of melting, overheating, cooling and casting of tensile

specimens is depicted in Figure 2.

Time

Tem

pera

ture

780 Co

980 Co

To castNo

overheating

After overheating

Heating

Cooling

880 Co

Fig. 2. Schematic representation of time-temperature treatment.

Mould QC 4080 PT-600-PvG

furnace

Crystaldigraph

recorder

PC Computer


The tensile test was carried out on an Instron 4469 machine at a rate of 20 mm/min. Studies were

carried out on a normal-running fatigue testing machine, which was the mechanically driven

resonant pulsator - Figure 3a. Specimens prepared according to PN-EN 1706 are shown in Figure

3b.

Fig. 3. a) Instron 4469 machine, b) samples of tensile strength.

Results and discussion

Since studies described in this paper are of a preliminary character only, the AlSi17Cu5Mg

silumin was cast without modification and refining.

At the first stage of investigations, the main parameters of the descriptive statistics of the

AlSi17Cu5Mg alloy overheated to a preset temperature and cooled at a constant rate of about

2,5oC/s were determined.

Next, the tensile strength Rm was determined, along with the respective mean values, intervals of

confidence, stability ranges and the scatter of results measured with standard deviation (SD). The

results are graphically depicted in Figure 4.


Fig. 4. Plotted relationship for the examined silumin tensile strength

in function of overheating temperature: a) 780oC, b) 880

oC, c) 980

oC.

At the next stage of investigations, the value of the tensile stress „σ” was determined and, based

on relationship (1), the value of the survival probability (PS), called model parameter, was

calculated.

( )

−=

m

sVP

0

0 expσσ

where:

PS(V0) – survival probability,

σ – tensile stress, [MPa],

σ0 – tensile stress, for which 37% of samples exceed this value in terms of test features,

[MPa],

m – Weibull modulus

Based on these data, the boundary value of σ0 was calculated for Weibull distribution along with

the value of modulus „m” considered a measure of the A390.0 alloy performance reliability when

overheated from two temperatures. The results are shown in Table 2.

Table 2. Parameters of reliability of the A390.0 cast alloy.

Overheating temperature, [oC]

Reliability parameter

σ0, [MPa] Modulus „m”

No overheating 780 136,3 7,12

After

overheating

880 121,8 8,56

980 98,5 12,69

(1)


Based on these values, graphical relationships for the tensile stress „σ” [MPa] in function of the

survival probability (PS) were determined, as shown in Figure 5.

0,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

70 80 90 100 110 120 130 140 150 160 170 180

Tensile stress, MPa

Su

rviv

al

pro

pab

ilit

y,

p

Fig. 5. Survival probability in function of tensile stress;

a) 780oC, b) 880

oC, c) 980

oC overheating temperature.

Summary

As proved by the investigations, melting, overheating and cooling to the pouring temperature

change in a significant way the value of the tensile stress:

σ0 (780oC) = 136,3 MPa,

σ0 (880oC) = 121 MPa,

σ0 (980oC) = 98,5 MPa, consequently affecting also the value of Weibull modulus: m(780

oC)=7,12;

m(880oC)=8,56 and m(980

oC)=12,69.

Hence a conclusion follows that the degree of molten alloy overheating reduces the tensile

strength of alloy, reducing also the scatter of the obtained results. Making the stability range of the

resultant feature more narrow, and hence increasing the alloy technological and operational stability

reduces the coefficient of variability and the boundary value of „σ0” for Weibull distribution,

increasing at the same time the value of parameter „m”. It has been assumed that this coefficient,

responsible for the survival probability (PS), determines the material reliability which for an

engineer-designer may constitute an important advantage over other materials. The reduced scatter

of the tensile stress values „σ” was confirmed by the reduced slenderness in Figure 5.

References

[1] C.L. Xu, H.Y. Wang, C. Liu, Q.C. Jiang, Crystal Growth 291 (2006), p. 540.

[2] Weimin Wang, Xiufang Bian, Jingyu Qin, S.I. Syliusarenko, Metall. Mater. Trans. 31A

(2000), p. 2163.

[3] H.S. Kang, W.Y. Yoon, K.H. Kim, M.H. Kim, P.Yoon, Mater. Sci. Eng. A 404 (2005), p. 117.

[4] Y.T. Pei, J.Th. M. De Hosson, Acta Mater. 49 (2001), p. 561.

[5] K.F. Kobayashi and L.M. Hogan, J. Mater. Sci. 20 (1985), p. 1961.

a)

σσσσ0

b) c)


[6] Z. Górny and J. Sobczak, Non-ferrous metals based novel materiale in foundry practice,

Copyright by ZA-PIS, Cracov, 2005.

[7] G.K. Gavrilin, Plavlenie i kristallizatsiya metallov i splavov, (Rus.) Metallurgiya, 1996.

[8] J. Szymszal, J. Piątkowski, T. Mikuszewski and M. Maliński, Arch. Foundry Eng. 9, issue 3

(2009), p. 195.

[9] J. Piątkowski and J. Szymszal, Transactions on transport systems telematics and safety.

Published by SilesianUniversity of Technology (2011), p. 150.

[10] Peijie Li, V. I. Nikitin, E. G. Kandalova and K. V. Nikitin, Mat. Sc. and Eng. A 332, (2002),

p. 371.

[11] C.L. Xu, Q.C. Jiang, Mater. Sci. Eng. 437 (2006), p. 451.

[12] J. Piątkowski, Sol. State Phen. 176 (2011), p. 29.

[13] J. Piątkowski, Arch. Foundry Eng. 10, issue 2 (2010), p. 103.


Mechanism of grain refinement in Al after COT deformation

Kinga Rodak1,a, Jacek Pawlicki1,b

1 Silesian University of Technology, ul. Krasińskiego 8, Katowice 40-019, Poland


Keywords: severe plastic deformation, aluminium, fine-grained structure, TEM, STEM

Abstract. The microstructure of Al processed by compression with oscillatory torsion (COT)

method have been studied. This method was applied to refine the grain structure to ultrafine

dimension. The aim of the study was to examine how severe plastic deformation technique (COT) -

alter the microstructure. The second aim is to understand the mechanism of grain refinement. The

microstructure was characterized using transmission electron microscopy (TEM) and scanning

electron microscopy (SEM) equipped with electron back scattered diffraction (EBSD) facility.

1. Introduction

The research of the methods of grain refinement is carried out in parallel to the intensive research

on the structural changes occurring in the deformed materials. Some metallic materials which are

deformed by means of Severe Plastic Deformation are characterized by ultrafine-grained and

sometimes even nanograined size. The interest in the development of bulk nanostructure arises

because the use of different SPD technologies provide new opportunities for developing

ultragrained and nanograined structure in metals. One of such methods is the compression with

oscillatory torsion (COT). This method has become recognized mainly as a method that enables

deformation of the materials to values of large plastic deformations, therefore, it is possible to

obtain a refined structure [1,2]. The benefits from applying the COT method are visible mainly in

the aspect of formation of a particular type of a spatial configuration of defects. Compression with

oscillatory torsion is a method of plastic deformation in which the material is deformed as an effect

of a changing deformation path. The appliance allows for the following parameters to change:

the compression velocity v, (the velocity of the lower punch shift). The maximal value of

compression velocity is 0,6 [mm/s],

the torsional frequency f. The frequency of the lower punch oscillation is regulated by the

inverter ranging from 0 [Hz] to 1,8 [Hz],

the torsional angle amplitude α. The set points of the kinematic magnitudes enable the

change of the torsional angle ranging from 0 [º] to ±6 [o],

the absolute strain ∆h [mm].

Using the different deformation parameters caused the presence of different phenomena that were

controlling the microstructure. The growth of deformation (realized through the Hz increase) causes

a progress in the grain refinement. The conducted deformation when the torsional frequencies were

0,8 Hz and 1,6 Hz is the most beneficial for obtaining the most refined grain. Using a significantly

higher torsional frequency 1,8 Hz during the deformation caused considerable restrictions in the

grain-refining, because of the intensive recovery which begins to dominate over the deformation

process. It can be noticed that a relatively small change of the compression rate, had an impact on

considerably greater refining of the structure. The acceleration of the deformation process by

increasing the compression rate (from 0,015 mm/s to 0,04 mm/s), causes the delay of the structure

recovery process. Despite reducing the equivalent deformation, from the reason of increase

compression rate, the progress in the structure refining is observed what denies the results described

in the literature [3,4]. However, the next increase of the compression rate ( 0,1 mm/s and 0,6 mm/s)

does not foster grain-refining. It should be explained by the fact that the equivalent deformation is

too small. The deformation is very dynamic what can indicate that there is not enough time for the

structure refining. The main effect of deformation is the increase of the structure’s homogeneity.

The homogeneity is obtained mainly by the increase of the absolute strain ∆h.


When the process parameters are as follows: the torsional frequency (0,8 Hz and 1,6 Hz), absolute

strain ∆h=7mm, and the compression rate (0,015 and 0,04 mm/s); the maximal refining of the Al

grain is obtained:

- the average diameter of the grain/subgrain correspondingly 600 nm and 300 nm,

- the fraction of high-angle boundaries is about 50% and

- the area fraction of the ultrafine grains is about 45%.

The aim of the present work was to describe of the mechanism of grain refinement that occur in Al

after COT deformation.

2. Experimental details

The tests were conducted on the samples from aluminium A0 (chemical composition is shown in

the Table.1). The samples for the tests were taken from the bars having 12 mm in diameter and then,

they were exposed to the heat treatment which involves the annealing in temperature of 200oC /1

hours. After this treatment the average diameter of the grain equaled 75 µm. The heating treatment

that was carried out, allowed for eliminating structural effects resulting from the previous

technological treatments and for obtaining the homogenous grain structure in the whole volume of

the material.

Table1. Chemical composition Al used in experimental

Chemical composition, wt [%]

Cu Fe Mn Mg Si Al

0.002 0.28 0.001 0.001 0.09 Ball.

The compression with oscillatory torsion method is regarded as a method characterized by the

heterogeneity of deformation. The most intense deformations proceed in places that are the nearest

to the lateral surfaces of the material, which is results from the functioning of the torsional moment.

The heterogeneity of the plastic deformation in the sample, causes the occurrence of a considerable

differentiation of the structure in its sectional view. Because of the heterogeneity of the

deformation, the microscope observations were carried out in areas located in a distance of about

0,8 of the sample’s radius. More information about dimension of sample used in experiment is seen

in [5].

The analysis of the dislocating structure was carried out using the Scanning Transmission Electron

Microscopy (STEM) technique which was applied thanks to the microscope Hitachi HD 2100A

equipped with the FEG gun, working at the accelerating voltage of 200 kV. With the help of

Transmission Electron Microscopy (TEM) Jeol 100B, the orientation of the grains were determined

based on the received pictures of Kikuchi lines. For the calculations the KILIN programme was

used that was developed on the University of Science and Technology in Cracow [6].

The detailed quantities research of the ultrafine-grained structures was conducted using Scanning

Electron Microscope (SEM) INSPECT F produced by FEI equipped with the gun with cold field

emission and the detector of electron back scattering diffraction (EBSD). In order to release the

structure of the material by using the SEM/EBSD method, firstly, the mechanical polishing was

used and then, electrolytic. The boundary between the grain and subgrain was determined on the

basis of the misorientation angle measurement. The divisional boundary was an angle equaling 15o.


3. Results of investigations

The deformed Al with low torsional frequency - 0,2 Hz is characterized mainly by the boundaries

with a small misorientation angle. The boundaries like HABs are seen in the fragmentary outline

(Fig.1). The dislocation cell structure and the DDWs dislocation walls with a high density of

dislocation are observed. The effects of arranging the dislocation structure are seen between the

particular DDWs (Fig.2).

Fig.1. EBSD maps illustrating Al microstructure

changes after COT processing: f=0,2 Hz;

α =±6o, v=0,015 mm/s and ∆h=7mm. The

HABs and LABs boundaries are respectively

shown as thick and thin lines

Fig.2. Al microstructure after COT deformation

with parameter: f=0,2 Hz, v = 0,015 mm/s, α

=±6o, ∆h=7 mm. DDWs boundaries are marked

with arrows

The large effective deformation (εf) realized by the increase of the torsional frequency to 0,8 Hz,

have an impact on the increase of the misorientation value between the created grains (Figs.3,4).

The great part of the analyzed surfaces are the banding structures isolated by high-angle boundaries

and elongated in accordance with the direction of the compression (Fig.3). STEM micrographs

(Fig.4) evidently demonstrate that deformation at higher value of εf parameter leads to generating

banded structure with low angle grain boundaries (Fig.4).

The map obtained using the EBSD technique show that the structure during the deformation process

where the torsional frequency was 1,6 Hz, is characterized by the considerable grain refining. The

deformation allows for the formation of the equiaxed structures. It was observed that the numerous

grains had the size no bigger than 1 µm (Fig.5). The increase of the effective deformation has an

influence on the generation of the HABs boundaries which are subject to the mutual intersection

(Fig.6). The result of this is the generation of the greater amount of the grains.

The Al grain/subgrain structures are complex. The obserwations indicate that the grain interiots can

be free from dislocations or have chaotically distributed dislocation. The small grains have sharp

grain boundaries and are almost free of dislocation. The dislocation cell structures were found

inside larger grains. The larger subgrains contains a high density of dislocation.

DDW

εf=15.4

εf=15.4








with parameters: f=0,8 Hz, v = 0,015 mm/s, α

=±6o, ∆h=7 mm. Subgrains are characterized by

misorientation approaching 10o. Visible

numbers corresponds to values of individual

areas misorientation







with parameters: f=1,6 Hz, v = 0,015 mm/s, α

=±6o, ∆h=7 mm. Arrangement of reciprocal

intersecting of LBs boundaries with HABs

misorientation type creates grain structure.

Visible numbers corresponds to values of

individual areas misorientation

200 nm

4,54o

7,11o

5,65o

10,23o

500 nm

30,03o

45,69o

28,03o

εf=61

εf=61

εf=120

εf=120

4,54o


Fig.7. Al microstructure after COT deformation with parameters: f=0,8 Hz, v = 0,015 mm/s, α

=±6o, ∆h=7 mm. ABC marked the appearing boundaries after changes in tilting angles: a) α= -13

o,

b) α=-7,9o. The diffraction patterns (axis zone [011]) recorded from the X area

Fig. 8. Al microstructure after COT deformation with parameters: f=1,6 Hz, v = 0,04 mm/s, α =±6o

,

∆h=7 mm; a) High misorientation of nonequilibrium grain boundaries in grain marked 1. Created

grain is characterized by various crystallographic orientation. Nonequilibrium grain boundary is

marked with arrow; b) Kikuchy diffractions with solutions. Numbers 1-4 in Fig.8a correspond to

Kikuchy patterns 1-4 for grains/subgrains; c) orientation of analized area

1

2

3

4

4

3 2 1

A B

X

A B

X

b) c)

a) b)

3

21

4

53,13o

15,52o

51,91o

3

21

4

53,13o

15,52o

51,91o

a)


The high-angle boundaries (HABs) marked in the Fig.7 have bulges characteristic for the

continuous dynamic recrystallization (CDRX) [7,8]. The sequence of figures registered during the

rotation of the sample at a given angle indicates that the bulges of HABs boundaries are not the

effects of the boundary migration but of the mutual superimposing of the boundaries which are in

one crystallographic orientation in a given microarea (Fig.7).This means that the creation of

ultrafine-grained structures using the COT method is not determined by the process of

recrystallization.

In order to trace the way in which the low-angle boundaries (LABs) change into HABs

boundaries, a series of structural tests was carried out in which the Kikuch diffraction was used.

Defining of the local orientations and crystallographic misorientations of particular areas,

allowed to formulate the mechanism in which the high-angle boundaries are formed. The examples

presented in Fig.8 suggest that the recovery of Al happens relatively fast. That is why the

dislocations generated during the deformation do easily give into annihilation. The orginal grains

are subdivided into smaller granular structure with high angles of misorientation and well- define

boundaries. The clusters of dislocation are usually created near the boundaries what causes

nonequilibrium state of the grains boundaries (Fig.8).

4. Discussion

A lot of place in the structural investigations is devoted to the mechanisms of grain-refining after

SPD deformations. This matter is interesting particularly because of the application of different

materials and different SPD techniques. The most important conclusions taken from the researches

on the mechanisms of grain-refining are as follows [9,10]:

the fragmentation of grain takes place when the dislocation boundaries taking different

forms are generated,

the deformation is accompanied by the processes of dynamic recovery or even

recrystallization.

An example of a material in which the grain-refining is the effect of deformation and of the

„extended recovery” is aluminum. The increase of misorientation happening thanks to the rotation

of the grains boundaries. The mechanism based on annihilation and absorption of the dislocation

through the grains boundaries.

Hughnes and Hansen [11] presented the concept of microstructure evolution for the classic

techniques of deformation which is based on the generation of dislocation boundaries which lead to

the division of the initial grains into smaller volumes. The proposed conception of the structure

evolution is also characteristic for SPD techniques because many researchers, introduce to the

description of the structure the terminology in a form of the shear bands or dislocation layers.

In general, there are three main mechanisms of the material structure-refining that are known [3,12]:

the production of new grains takes place thanks to a gradual increase in misorientation of

dislocation boundaries as a result of absorption of new dislocations created during the

deformation,

the fragmentation of grains takes place thanks to the generation of the shear bands,

the fragmentation of grains takes place thanks to the production of new grains as a result of

the continuous recovery or continuous recrystallization.

The refining of the Al grain after the COT deformation, happens as a result of the generation of

dislocation boundaries, which together with the growth of deformation transform themselves into

ultrafine-grained structure. The introductory stage of the grain-refining is the creation of DDWs

dislocation walls, the misorientation of which reaches even a few degrees and they stretch along the

considerable fraction of a grain, separating blocks of dislocation cells (CBs) (Fig.2). Within the

boundaries a high density of dislocation is accumulated.


The growth of deformation causes the transformation of the DDWs dislocation walls into lamellar

boundaries (LBs) which has larger misorientation (sometimes above 15o - HABs) which resembles

long subgrains (Fig.3). Inside of the lamellar boundaries, the dislocation structure is regular and

singular dislocation cells are usually noticeable (Fig.4). Moreover, the distance between the lamellar

areas decreases (Fig.1 and Fig. 3, Fig.2 and Fig.4). It can be assumed that the accumulation of

deformation induces not only the creation of new dislocation boundaries and high angle boundaries

but above all, it induces the crossing of dislocation boundaries. This phenomenon of intensive

boundaries crossing (Figs.5,6) is a result of activating the subsequent slip systems. The places

where the dislocation boundaries are crossing induce the generation of almost equiaxed

subgrains/grains (Fig.6). The result of EBSD test also constitute the confirmation of the STEM

investigations. On the basis of EBSD investigations it was proved that in a great deal of cases

elongated neighboring grains remain in crystallographic compatibility. In the case of fine, equiaxed

grains, the orientation was accidental (Figs.3,5). The structural analysis presented here show that

the dominating mechanism of forming the high-angle grains in Al is the growth of misorientation in

dislocation boundaries. This, in turn, happens as an effect of absorbing dislocations to the grains

boundaries during the deformation process.

5. Summary

The process of grain refining proceeds by the generation of the LABs and HABs dislocation

boundaries. In the introductory stage of deformation the dislocation boundaries are formed which

are perpendicular to the direction of the compression force. The formation of the dislocation

boundaries which proceed in such a way, suggests that in the initial stage of deformation it is mainly

the compression that initiates the process of the grain refining. The non-directional process of

deformation (the introduction of an additional torsion causing the change in the direction of

loading) leads to the deformation of the material in more and more numerous systems of glides. The

effect of the introduced, additional loading is the increase in the number of the dislocation

boundaries that cross mutually. When the effective deformation in the microstructure increases, the

distances between the dislocation boundaries decrease – a new order of LBs dislocation boundaries

are created. A significant role in forming the ultrafine-grained structure has the recovery process.

Dislocations are rearranged, undergo annihilation and are also absorbed to the grain boundaries.

Such a rebuilding of a dislocation structure causes the increase of the misorientation within the

grain boundaries.

References

[1] G. Niewielski, D. Kuc, K. Rodak, F. Grosman, J. Pawlicki, Influence of strain on the copper

structure under controlled deformation path conditions, Journal of Achievments in Materials

and Manufacturing Engineering, 17 (2006) 109-112.

[2] K. Rodak, J. Pawlicki, Microstructure of ultrafine-grained Al produced by severe plastic

deformation, Archives of Materials Science and Engineering, 28 (2007) 385-448.

[3] M. Richert, Effect of large deformations on the microstructure of aluminium alloys, Materials

Chemistry and Physics, 81, (2003) 528-530.

[4] H. Petryk, S. Stupkiewicz, A quantitative model of grain refinement and strain hardening

during severe plastic deformation, Materials Science Engineering, A 444, (2007) 214-219.

[5] K. Rodak, J. Pawlicki, Microstructure of ultrafine-grained Al produced by severe plastic

deformation, Archives of Materials Science and Engineering, 28 (2007) 385-448.

[6] M. Richert, K. Chruściel, J. Długopolski, A. Baczmański, Computer program „Kilin”.

Orientation, desorientation and microtexture.

[7] H. Hollberg et al., Modeling of continous dynamic recrystallization in commercial-purity

aluminium, Materials Science and Engineering A 572 (2010) 1126-1134.


[8] M. Eizadjou et al., Microstructure and mechanical properties of ultra-fine frains (UFGs)

aluminium strips produced by ARB process, Journal of Alloys and Compounds, 474 (2009)

406-415.

[9] Y.T. Zhu, Performance and application of nanostructured materials produced by severe plastic

deformation, Sctipta Materialia, 51 (2004) 825-830.

[10] R.Z. Valiev, Producing nanostructured materials by severe plastic deformation for advanced

applications, Metting Procedings RTO-MP-AVT-122, 1-12.

[11] D.A. Hughes, Microstructure evolution, slip patterns and flow stress, Materials Science

Engineering, A319 (2001) 46-54.

[12] M. Lewandowska, Microstructure and properties of aluminium alloys processed by

hydrostatic extrusion, Oficyna Wydawnicza Politechniki Warszawskiej, 2006 [in polish].


Deformation-induced grain refinement in AlMg5 alloy

Kinga Rodak1,a, Jacek Pawlicki1,b, Marek Tkocz1,c

1 Silesian University of Technology, ul. Krasińskiego 8, 40-019, Katowice, Poland

a [email protected], b [email protected], c [email protected]

Keywords: severe plastic deformation, aluminium alloy, fine-grained microstructure, MAXStrain, mechanical properties

Abstract. The results presented in this paper are concerned with the microstructure and the

mechanical properties of the AlMg5 alloy subjected to severe plastic deformation by multiple

compression in two orthogonal directions. Four experiments with an increasing number of passes

were conducted on the Gleeble MAXStrain system in order to obtain various effective strain levels.

The microstructure of the most deformed, central parts of samples was investigated by means of the

light microscopy (LM) and the scanning transmission electron microscopy (STEM). The

mechanical properties of the analyzed sample regions were determined as well. Investigations

revealed that severe cold deformation of the AlMg5 alloy leads to strong grain refinement.

Moreover, fragmentation of large intermetallic inclusions and their regular distribution were

obtained. Microstructural changes led to significant improvement in the strength properties. After

reaching the effective strain of 9, the AlMg5 alloy exhibited UTS, YS and HV values almost two

times higher than corresponding values determined for the starting, annealed material.

Introduction

The consequence of severe plastic deformation (SPD) is the crystal fragmentation of a large

number of metallic alloys to an ultrafine or even nano-grained dimension. SPD can be performed by

various methods reported in literature, e.g. equal channel angular pressing (ECAP), hydrostatic

extrusion (HE), high pressure torsion (HPT), compression with reversible torsion and multiple,

cyclic compression in the orthogonal direction [1-4]. Thanks to the refinement of microstructure,

high strength is commonly achieved in the light alloys, e.g. Al-Mg alloys [5-7]. Besides of the low

specific weight, the aluminum alloys with magnesium are characterized by high resistance to

corrosion [8,9]. Therefore, obtaining the significant improvement in the strength properties makes

these materials an attractive alternative for many applications.

The present work was aimed to describe the microstructure and the intermetallic inclusions

evolution in AlMg5 alloy during cold severe plastic deformation imposed by consecutive two-axial,

multiple compression and how these microstructural changes affect the mechanical properties of the

material investigated.

Material and methods

Starting material. Investigations were performed on AlMg5 (5019) alloy, the chemical

composition of which is given in Table 1. The cold drawn bars were homogenized at 500oC for 2h

and then cooled slowly down, to obtain the grain size of about 50 µm. Metallographic observations

revealed that the alloy microstructure consists of the matrix and the intermetallic inclusions of

various sizes. The inclusions are arranged in chains along the direction of drawing and occur in

clusters (Fig.1).

Table 1. Chemical composition of the investigated AlMg5 alloy, % wt.

Mg Mn Fe Si Zn Cr Cu Ti Ni Al

5.1 0.35 0.5 0.4 0.02 0.2 0.1 0.2 0.03 Balance


Fig. 1. Optical micrograph of the starting alloy after drawing and annealing at 500oC for 2h

Cold forming. The prepared starting material was then subjected to severe cold plastic

deformation by means of the Gleeble MAXStrain system. Samples with the square cross section

(10×10 mm) and the length of 27 mm, unrestrained lengthwise, were subjected to multiple

compression with 90 deg rotation between subsequent passes. Four different tests were conducted

with different number of passes (4, 8, 16 and 32). Anvils of 10 mm in length were used. The

samples were deformed with the average strain rate of 0.5 s-1

. According to the experimental data

acquired, the actual true strain values (calculated as the natural logarithm of the initial to final

sample height ratio) in the subsequent passes varied within the range of 0.15 - 0.3.

Microstructure investigations. HITACHI HD-2300A STEM microscope with the field

emission type gun was applied for microstructure investigations. In this study, the central parts of

compressed samples were taken into consideration. Foils for the STEM examinations were

electropolished using the A2 electrolyte. The light microscopy was also carried out to reveal the

intermetallic inclusions and their distribution in the alloy matrix. Metallographic sections were cut

along the longitudinal axis of the compressed samples.

FEM simulations. Due to the presence of friction and material outside the deformation gap,

distribution of the local effective strain is obtained in the sample after compression. The highest

strain accumulation occurs in the central part of a sample (Fig. 2). The effective strain in this region

is greater than the true strain determined from the sample height reduction. In order to evaluate the

actual total values of the effective strain in a sample centre after subsequent passes, numerical

simulations corresponding to the experiment conditions were executed. Forge2009 FEM software

were applied in this study.

Microhardness measurements. The microhardness was evaluated by the Vickers method at

loads of 200 g. The Future Tech FM 700 unit was applied in this case.

Tensile tests. Determination of the mechanical properties such as the yield strength (YS), the

ultimate tensile strength (UTS) and the uniform elongation (uEL) was performed on the MTS

QTest/10 machine equipped with digital image correlation system (DIC) [10]. The DIC method is

based on computational algorithms that track the grey value patterns in digital images of test

surfaces, taken before and after an event that produces surface displacements. The precision of the

method is of the order of 2/100 pixel while the minimum detectable displacement is of the order of

1/100 pixel. The minisamples with the total length of 2.2 mm (Fig. 3) were cut out of the central

parts of samples subjected to multiple compression in order to get the results representative to the

microstructures analysed.


Fig. 2. The shape and the effective strain

distribution on a half of the compressed

sample after 32 passes (the central cross

section is seen at the bottom)

Fig. 3. The shape and dimensions of

minisamples, cut from the centre of

compressed specimens and used in the static

tensile tests for determining the mechanical

properties


Effective strain evaluation. Numerical simulations allowed for estimation of the actual, total

effective strain values in the central parts of samples subjected to multiple compression. These

values - for the conducted and mentioned earlier deformation variants - are collected in Table 2.

Table 2. Estimated total values of the effective strain at the centers of compressed samples

Number of passes 4 8 16 32

Total effective strain 2.1 3.8 6.2 9.1

STEM examinations. Results of the STEM investigations indicate that after obtaining the

effective strain of 2.1 in the central area of a sample, fragmentation of the original grains occurs. As

expected, multiple compression in two orthogonal directions results in development of the

characteristic deformation bands (DBs) (Fig. 4a) and dislocation tangles (Fig. 4b).

After reaching the effective strain of 3.8, the new microstructural component has been observed.

The crossing of deformation bands form a structure which is similar to a dislocation cell structure.

This indicates that new slip systems operated during deformation (Fig. 4c). A new local high-angle

boundary grain appeared with the diameter of 300 nm (Fig. 4d).

As the effective strain increased to 6.2, another new local high-angle boundary grains appeared

(Fig. 4e-f), but the microstructure in many areas is composed of weakly distinguishable subgrains

outlined by the low-angle dislocation subboundaries. The dislocation substructure appears as fine

irregular fragments. Generally, the microstructure contains the ultrafine subgrains with a high

density of dislocations (Fig. 4f).


Fig. 4. AlMg5 alloy microstructure at the specimen’s centre after compression in the MaxStrain

system, a-b) ε=2.1, c-d) ε=3.8, e- f) ε=6.2, g-h) ε=9.1

a) b)

c) d)

e) f)

g) h)


The presence of small equiaxed grains with the sharp grain boundaries that contain a high

density of dislocations was observed in the most deformed sample. STEM provided that applied

multiple compression was sufficient to deform the microstructure at the samples centers uniformly

and to produce grains having a high fraction of the high-angle boundaries (Fig. 4g-h). The presence

of Mg atoms reduces the dislocation mobility and introduces solid solution strengthening, and so

does the rate of recovery in Al-Mg alloys. For this reason many of the boundaries are not well

defined. This feature was attributed to development of the arrays of high- energy noequilibrium.

Analysis of inclusions. One of the characteristic features of the AlMg5 alloy microstructure is

the presence of the intermetallic inclusions (e.g. with Fe, Si, Mn and Cu) which form during

solidification and are relatively large, although their size remains at the micrometer level.

The light microscopy (LM) observations of inclusions revealed that they significantly changed

size, shape and distribution after deformation (Fig. 5). It was found that multiple compression

conducted by the MAXstrain system causes fragmentation of the intermetallic inclusions. The

higher strain, the larger fragmentation and more homogenous distribution of inclusions are

obtained. However, some large particles still remains after reaching the effective strain of about 9

(Fig. 5c).

Fig. 5. LM micrographs of the AlMg5 samples’ centres before (a) and after compression in the

MAXStrain system: b) ε=3.8, c) ε= 9.1

a)

b) c)


The examples of fragmented inclusions, noticed after multiple compression, are presented in

Figs. 6 and 7.The EDS mapping of the observed inclusions distinguished three kinds of particles in

the alloy, containing mainly: Si (Fig. 6c and Fig. 7b), Mn (Fig. 6d), Fe (Fig. 6f) and Cu (Fig. 6e).

Fig. 6. Examples of the fragmented inclusions in the AlMg5 alloy microstructure after reaching

the effective strain of 6.2: a) secondary electron image, b) phase contrast image, c) X- ray mapping

of Si, d) X- ray mapping of Mn, e) X- ray mapping of Cu, f) X- ray mapping of Fe

b)

Si c) d) Mn

e) f) Cu Fe

a)

fragmented

inclusions


Fig. 7. The partly crushed inclusion in the AlMg5 alloy microstructure after reaching the effective

strain of 9.1: a) secondary electron image, b) X- ray mapping of Si, c) X- ray mapping of Al,

d) X- ray mapping of Mg

Mechanical properties. Values of the yield stress (YS), the ultimate tensile strength (UTS), the

uniform elongation (uEL) and the hardness (HV) are collected in Table 3. It should be noted that the

mechanical properties are related only to the central parts of compressed samples. Therefore, they

are representative to the microstructures presented in this paper and correspond to the total values of

the effective strain evaluated. It is clearly seen that multiple compression caused significant strain

hardening of the alloy. It’s obvious that the uniform elongation decreases with the rise of the

effective strain, however it still remains on a satisfactory level.

Table 3. The mechanical properties at the central regions of the AlMg5 alloy samples

Effective

strain

Hardness

[HV]

UTS

[MPa]

YS

[MPa]

uEL

[%]

0 (initial state) 77 234 143 18.2

2.1 118 310 276 15.2

3.8 122 - - -

6.2 123 384 326 10.8

9.1 141 416 341 10.7

a) b)

c) d)

Si

Al Mg

partly

crushed

inclusion


Conclusions

The results of presented investigations allow to draw the following conclusions:

1. Application of the multiple compression in two orthogonal directions involves a considerable

refinement of the AlMg5 alloy microstructure. The method is effective in generating HABs.

2. The applied forming technique also results in fragmentation of the intermetallic inclusions.

After severe plastic deformation the inclusions are more regularly distributed in the matrix.

3. Evolution of the AlMg5 alloy microstructure in samples subjected to severe plastic deformation

indicates a transition from the microstructure dominated by cells with low angle boundaries to

the equiaxed nano- and ultra-grained microstructure.

4. The multiple compression results in the significant increase of the strength properties in

comparison with the corresponding properties of the investigated alloy in the initial, annealed

state.

References

[1] K.J. Kurzydłowski, Microstructural refinement and properties of metals processed by severe

plastic deformation, Bulletin of the Polish Academy of Sciences, Technical sciences, 52, 4 (2004)

301-311.

[2] K. Rodak, J. Pawlicki, Effect of compression with oscillatory torsion processing on structure

and properties of Cu. J. Mat. Sci. Technol., 27 (11) (2011) 1083-1088.

[3] H. Petryk, S. Stupkiewicz, A quantitative model of grain refinement and strain hardening

during severe plastic deformation, Mat. Sci. Eng. A 444 (2007) 214-219.

[4] R. Kuziak et al., New possibilities of achieving ultrafine grained microstructure in metals and

alloys employing MaxStrain technology, Solid State Phenom., 101-102 (2005) 43-48.

[5] M. V. Markushev et al., Structure and mechanical behavior of an Al-Mg alloy after equal

channel angular extrusion, NanoStructured Materials, 12 (1999) 839-842.

[6] M. Richert et al., Work hardening and microstructure of AlMg5 after severe plastic

deformation by cyclic extrusion and compression, Mat. Sci. Eng. A355 (2003) 180-185.

[7] T. Kovarik et al., Mechanical properties and microstructure evolution in ECAP processed Al-

Mg-Si alloy by ECAP deformation, Proc. of the Int. Conf. Metal 2010, 18-20.05.2010, Roznov pod

Radhostem (2010)

[8] G. Nurislamova et al., Nanostructure and related mechanical properties of an Al-Mg-Si alloy

processed by severe plastic deformation, Philosophical Magazine Letters, 88 (2008) 459-466.

[9] M. Greger et al., The formation of submicron and nanocrystaline grain structures in Al-Mg-Si

alloy, Proc. of the Int. Conf. Metal 2010, 18-20.05.2010, Roznov pod Radhostem (2010)

[10] L. Chevalier, S. Calloch, F. Hild, Y. Marco, Digital image correlation used to analyze the

multiaxial behavior of rubber –like materials, Eur. J. Mech. A Solids, 20 (2001) 169-187.


CMT and MIG-Pulse robotized welding of thin-walled elements made of 6xxx and 2xxx series aluminium alloys

Janusz Adamiec1,a, Tomasz Pfeifer2,b, Janusz Rykała2,c

1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

2 Institute of Welding, ul. Błogosławionego Czesława 16-18, 44-100 Gliwice, Poland


Keywords: aluminium alloys, CMT, low energy welding, microstructure analysis, heat treatment

Abstract

The article presents the course and the results of research on material and technological welding

conditions of different aluminium alloys using standard (MIG-Pulse) and low energy welding

method (CMT) as well as discusses the properties of welded joints and the application fields of

modern low energy welding devices for joining thin aluminium sheets.

Introduction

The necessity to build structures of lower weight, yet characterised by adequate strength is the

reason for an increase in the application of high-strength aluminium alloys in various industries. In

consequence, there is also a growing demand for welded joints of these alloys, characterised by

appropriate quality and mechanical properties [1-5].

For a number of reasons, so far the application of the MIG method in welding of thin-walled

elements made of high-strength plastic-worked and precipitation-hardened aluminium alloys has not

proved fully satisfactory; this being in particular due to significant porosity (esp. 2xxx-series alloys)

and too much heat supplied to metal, which, in turn, resulted in a high decrease in mechanical

properties in the fusion zone (caused by a loss of advantageous output structure of alloy following

precipitation hardening) as well as susceptibility to hot cracking (formation of low-melting

eutectics). In addition to the foregoing, traditional MIG welding is accompanied by significant

spatter and deformation of elements being joined, which decreases the aesthetics of finished

products and requires time-consuming post-weld treatment or the application of procedures

preventing the aforesaid phenomena, which, in turn, is responsible for lower efficiency and

complicated design of instrumentation and technological processes [2-7].

The recent years have seen research and development carried out by leading manufacturers of

welding equipment aimed to develop low-energy MIG/MAG (CMT, ColdArc etc.) welding, having

in view solutions to problems which accompany the joining of thin-walled elements characterised

by limited weldability, post-weld porosity and sensitivity to heat effect [1-2]. Publication [1]

presents the principle of operation of these methods and initial results of tests dedicated to their

application in the welding of aluminium, including technological issues and the influence of a given

method on the aesthetics of joints. This article presents the impact of a low-energy CMT method on

the structure of welded joints made of Al-Mg-Si and Al-Cu aluminium alloys as well as the

former’s susceptibility to cracking.

Course and results of tests

The technological tests of a welding process by means of CMT and MIG-Pulse methods

incorporated the production of butt joints using the following combinations of base and filler metals

(base metals designated acc. to PN-EN 573-3, filler metals acc. to PN EN ISO 18273) :

• 2.0 mm-thick Al-Mg-Si plate, grade EN AW 6082 and electrode wire AlMg4.5MnZr (Al 5087)

φ1.2 mm,


• 2.0 mm-thick Al-Cu plate, grade EN AW 2017A and electrode wire AlCu6MnZrTi (Al

ML2319) φ1.2 mm.

As opposed to the most commonly applied filler metals such as Al-Si (AlSi5 and AlSi12) or Al-

Mg (AlMg5Cr), the foregoing are characterised by slightly inferior plastic properties but superior

mechanical properties, which ensure good metallurgical properties of the weld and allow post-weld

precipitation hardening of joints made of the aforesaid aluminium alloys [8-9]. Both filler metals

contain zirconium stabilising the structure and preventing grain growth. Butt joints were produced

for each of the combinations (i.e. base metal-filler metal) presented above. Due to the application of

specialist consumables (characterised by complex composition) and lack (in most cases) of software

related to the aforesaid welds, it was necessary to take advantage of current characteristics

developed for electrode wires Al-Mg or Al-Si by welding equipment manufacturers.

Technological tests involving both methods in question were carried out on a station provided

with a Fronius-manufactured device TransPuls Synergic 2700 (CMT method) and a Cloos-made

device GLC 553 MC3R (MIG-Pulse method) combined with Cloos’s robot ROMAT 310, thus

ensuring repeatable conditions of tests. The technological tests of welding the aforementioned

alloys by means of the CMT and MIG-Pulse methods revealed that there is a possibility of applying

such welding parameters which make it possible to obtain a welded joint characterised by the

quality level B following the requirements of the standard PN EN ISO 10042.

In addition to the technological tests mentioned above, additional welding trials of single-sided

stiffened joints were carried out in order to measure and compare the angular deformation α of the

joints made by means of both methods (Fig. 1). For each combination (i.e. base metal-filler metal) 5

test joints were produced. The results of measurements were averaged and presented in Table 1.

Fig. 1. Angular deformation of joint

Table 1. Measurements of angular deformations α in aluminium alloy joints

made with CMT and MIG-Pulse methods

Alloy grade

(welding method)

Deformation angle

α* [ o]

EN AW 6082

(MIG-Pulse) 7,5

EN AW 6082

(CMT) 5,5

EN AW 2017A

(MIG-Pulse) 7,5

EN AW 2017A

(CMT) 4,5

* - average value of 5 measurements of

welded joints

The research also involved hot crack tests utilising the Houldcroft method, carried out on test

pieces made of 4.0 mm-thick EN-AW 6082 and EN-AW 2017A aluminium plates, welded using the

CMT and MIG-Pulse methods. Each of the methods was used to perform three overlay welding


tests on previously prepared plates. The aforesaid tests were conducted ensuring such conditions

and parameters as were necessary to obtain proper fusion into the base metal (a slight bulge of

metal visible on the test piece from the root side).

Three of the test pieces made of the EN-AW 6082 alloy applying the CMT method revealed no

cracks. In turn, in case of the same alloy and the MIG-Pulse method some slight cracks were

revealed from the “root side”. The two remaining test pieces were free from any cracks. The

samples made of the EN-AW 2017A alloy, overlay-welded with the CMT method, did not reveal

any surface imperfections, whereas the test pieces made of EN-AW 2017A alloy, subjected to

overlay welding with the MIG-Pulse method cracked right through (Fig. 2 and 3), both in the weld

itself and in the zone near the weld.

Fig. 2. Result of Houldcroft test made with MIG-Pulse method, test piece made of EN-AW

2017A alloy

Fig. 3. Result of Houldcroft test made with CMT method, test piece made

of EN-AW 2017A alloy

The next stage of the tests consisted in the selection of some of the joints made of the 2.0 mm-

thick plates and subjecting them to heat treatment (supersaturation followed by artificial ageing).

Afterwards, the test pieced (in the state preceding and following the heat treatment) underwent

metallographic examination. All test pieces were etched in Keller reagent. Examples of

macroscopic metallographic photographs are presented on Fig. 4 and 5.


Fig. 4. Macrostructure of 2.0 mm-thick Al-Cu joint (grade EN AW 2017) welded with CMT

method, filler metal AlCu6MnZrTi φ1.2 mm, mag. 3x.

Fig. 5. Macrostructure of 2.0 mm-thick Al-Cu joint (grade EN AW 2017) welded with CMT

method, after heat treatment, filler metal AlCu6MnZrTi φ1.2 mm, mag. 3x.

The selected test pieces were also subjected to microscopic metallographic examination (light

microscopy and scanning electron microscopy). Test pieces for light microscopy were etched using

Keller reagent and tests pieces for scanning electron microscopy were not etched. The results of the

examination are presented on fig. 6.

The aluminium alloy joints were also subjected to chemical composition analysis carried out

with a scanning microscope HITACHI S-4200, provided with X-ray microanalysis system NORAN

Voyager 3500 and an EDS spectrometer. The chemical composition tests were conducted at an

electron beam accelerating voltage of 15 keV. The local chemical composition microanalysis of the

weld was supplemented by surface distribution of chemical elements in the zone near the

weld/HAZ.

The structure of the weld of the joints made of EN AW 6082 alloy revealed the presence of

dendrites of the solid solution of magnesium in aluminium α-Al and fine precipitates of

intermetallic phases on the boundaries of these dendrites [9]. The analysis of the metallographic

examination results of the test pieces prior to heat treatment revealed the presence of bright globular

intermetallic phases containing Si (approx. 3% by weight), Mn (approx. 5.5% - 7.2% by weight)

and iron ((approx. 10% by weight). In combination with aluminium the above chemical elements

can form the phases Al3Mg2, Al3Fe and AlMg2Mn. In turn, dark globular phases contain Mg and Si

(probably forming a phase Mg2Si). After heat treatment, the test specimens revealed fine

precipitates of the bright phase on the boundaries of the dendrites (rich in Mg, Si, Mn and Fe) as

well as dark phases situated also on the boundaries of the dendrites, in the form of discontinuous

lattices, containing mainly Mg and Si. The latter form, as in case of the state preceding the heat

treatment, phases Mg2Si, yet of different morphology [9].

The weld of the joints made of EN AW 2017 A alloy, not subjected to heat treatment, revealed

the presence of dendrites of the solid solution of copper in aluminium and globular precipitates of

the bright phase on the boundaries of crystals (Fig. 6c-e) [7]. The phase in question contains approx.

40-45% Cu, which may indicate that it is Al2Cu (Fig. 8). The latter is confirmed by the surface

distribution of chemical elements including copper and aluminium in the phase (Fig 9) [7]. The

structure analysis at greater magnification revealed a dark globular phase (Fig. 6f). The phase

contained Si and Mg (possible phase Mg2Si). Similar structures were revealed in the test pieces

following heat treatment (Fig. 7). The weld is composed of dendrites of the solid solution of copper

in aluminium and the precipitates of an intermetallic phase on the boundaries of the dendrites in the

form of a discontinuous lattice (Fig. 7c); the chemical composition corresponding to that of Al2Cu

(Fig. 10). In inter-dendritic areas on the boundary with the phase Al2Cu it was also possible to

observe a dark phase rich in silicon (approx. 6%) (Fig. 7d,e; Fig. 10) [7].


Base metal; magn. 100x

HAZ; magn. 100x

Fig. 6. Microstructures of 2.0 mm-thick EN AW 2017A alloy made with CMT method; fusion zone

view (left side – weld, right side – base metal)

a) b)

c) d)

e) f)


Base metal; magn. 100x

HAZ; magn. 100x

Fig. 7. Microstructures of 2.0 mm-thick EN AW 2017A alloy made with CMT method, after heat

treatment; fusion zone view (left side – weld, right side – base metal)

a) b)

c) d)

e) f)


Weight % Al-K Cu-K

M4(4)_pt1 95.59 4.41

M4(4)_pt2 92.46 7.54

M4(4)_pt3 54.44 45.56

Atom % Al-K Cu-K

M4(4)_pt1 98.08 1.92

M4(4)_pt2 96.65 3.35

M4(4)_pt3 73.78 26.22

Fig. 8. Results of microanalysis of EDS chemical composition of area near weld, of 2017 A

aluminium alloy welded joint



aluminium alloy welded joint – surface distribution of chemical elements in weld


Weight % Al-K Si-K Cu-K

M40(4)_pt1 87.88 6.03 6.09

M40(4)_pt2 59.91 40.09

M40(4)_pt3 98.39 1.61

M40(4)_pt4 58.84 6.83 34.33

Atom % Al-K Si-K Cu-K

M40(4)_pt1 91.29 6.02 2.69

M40(4)_pt2 77.88 22.12

M40(4)_pt3 99.31 0.69

M40(4)_pt4 73.57 8.20 18.22


aluminium alloy welded joint – after heat treatment


The research also involved tensile tests (acc. to PN-EN ISO 4136) of the welded joints before

and after heat treatment (results marked in grey and with an asterisk), produced by means of both

methods (i.e. CMT and MIG-Pulse), as well as bend tests (PN-EN ISO 5173:2010). The results of

the tensile tests are presented in Table 2; the value of Rm(w) being the average of 3 tests. The said

table also presents the strength of the base metal following PN-EN 485-2:2007 and the minimum

required tensile strength of the joints welded following the requirements of PN-EN ISO 15614-2.

Table 2. Results of tensile tests, tensile strength of base metal and minimum required strength of

welded joint

Alloy grade

(welding method)

Rm(w)1)

[MPa]

Rm(pm)2)

[MPa]

Min. Rm(w)3)

[MPa]

EN AW 6082

(MIG-Pulse)

215,7

280 186 442,2*

EN AW 6082

(CMT)

232,3

468,6*

EN AW 2017A

(MIG-Pulse)

266,3

390 225 320,6*

EN AW 2017A

(CMT)

273,5

325,5*

Note:

1) Rm(w) – tenslile strength of welded joint,

2) Rm(pm) – tensile strength of base metal,

3) Min. Rm(w) – required minimum tensile strength of welded joint

acc. to PN-EN ISO 15614-2:2008,

* – results of welded joints after heat treatment.

Analysis of results and conclusions

The joints of both metal combinations (i.e. base metal-filler metal) produced by means of the

MIG-Pulse method were characterised by inferior aesthetics if compared to that of the joints made

using the CMT method. During a visual inspection no imperfections such as pores were detected,

yet the appearance of the face of the weld (particularly in case of the joints made of EN AW 2017A

alloy) indicated the presence of blowholes just underneath its surface, which was later confirmed by

radiographic tests. The joints produced with the CMT method did not contain or contained fewer

blowholes. The surface of the joints made with the Cloos-manufactured device was covered with a

layer of oxides (difficult to remove) and slight traces of spatters. In case of the CMT method no

spatters were detected and the layer of oxides either did not exist or was thinner and easy to remove.

This phenomenon can be attributed to the specific character of the CMT method and smaller

amount of supplied heat. Although it is either impossible or extremely difficult to precisely

determine the real value of linear energy in case of both of these methods, the practical proof of this

conclusion is easy to provide. Directly following the welding process, the joints made with the

CMT method could be taken off the stand bare-handed, unlike in case of the joints made with the

MIG-Pulse method. The post-weld temperature of the joints produced with the CMT method was

lower.

The observation is also confirmed by measurements of angular deformations in the single-sided

stiffened joints, which proved lower in case of the joints made with the CMT method; this being

due to a smaller amount of supplied heat and different shapes of welds in case of the methods under

test.


The macroscopic tests revealed that the joints produced with the CMT method, if compared to

those made with the MIG-Pulse method, were characterised by more uniform fusion and most

welds had regular, elliptical faces. The effect of a flat face and a bigger root in case of the welds

made with the MIG-Pulse method results from the specific character of the latter and, in particular,

from a higher arc voltage (and respectively lower current) in spite of comparable values of linear

energy (calculated in a classical manner) for these combinations of base/filler metals. The joints

made of EN AW 2017A alloy revealed greater porosity. Each of the joints produced with the CMT

method was characterised by very good aesthetics and an even face.

The surface of the joints did not reveal any hot cracks, typical of these base metals. Only in case

of the metallographic specimens of the joints made of EN AW 6082 alloy by means of the MIG-

Pulse method it was possible to detect micro-cracks along the grain boundary. The hot crack

resistance tests conducted using the Houldcroft method confirmed that the application of the CMT

method reduces the possibility of the formation of hot cracks in welded joints made of high-strength

aluminium alloys. Nonetheless, the reduction of hot cracks should, first of all, be attributed to

appropriate selection of filler metals, low linear energy and characteristics of both methods,

responsible for more precise transfer of liquid metal in the arc if compared to the classical MIG-

Pulse method.

Detailed metallographic tests made it possible to ascertain that the structure of aluminium alloy

joints is characteristic of the applied base metal-filler metal combinations and conducted heat

treatment. The analysis of chemical composition did not reveal differences as to the amount of

precipitates of chemical elements and phases in relation to the method by means of which a given

joint was made.

The tensile tests revealed that all of the welded joints met the minimum strength-related

requirements. Only in case of the joints made of EN AW 6082 alloy (produced by means of both

welding methods), it was possible to observe a rupture occurring in the base metal. In case of the

test pieces made of EN AW 2017A alloy, the rupture always occurred in the weld. The same

situation could be observed in case of the joints subjected to heat treatment. The results of the

tensile tests do not vary significantly (approx. 10-20 MPa) in case of both welding methods, yet in

most of the base metal-filler metal combinations, better Rm(w) results were obtained if the CMT

method was applied. The strength-related test results of the joints after heat treatment indicate its

proper course in case of EN AW 6082 alloy. In case of EN AW 2017A alloy, the strength-related

results (320-350 MPa) following heat treatment were not satisfactory and their value should be

significantly higher (over 400 MPa). The reason for such a state was probably too long hold time

during ageing. The aforesaid situation resulted in the excessive precipitation of a phase Al2Cu

(confirmed by further microscopic examination and an EDS phase analysis). The result of the

foregoing was “the overageing of the alloy” and decreased strength of joints. The bend tests of the

joints made of alloy EN AW 6082, preceding and following heat treatment, made by means of both

methods, produced positive results. In turn, in case of EN AW 2017A alloy all of the test pieces

were destroyed before reaching a required angle of 180°. Too low values of plastic properties can

probably be attributed to the intermetallic phases formed in the HAZ.

The analysis of the results obtained in the technological, strength-related and metallographic

tests made it possible to formulate the following conclusions:

1. The CMT and MIG-Pulse welding methods enable obtaining butt joints characterised by very

good quality and mechanical properties, allow the application of both alloys (plates with a

thickness of over 2 mm). It was not possible to obtain good plastic properties for the joints

made of EN AW 2017A alloy, which can probably be attributed to its limited weldability.

2. All of the joints made with the CMT method were characterised by superior aesthetics than

those produced with the MIG-Pulse method; aesthetic appearance being an important

evaluation factor in today’s industry.


3. The application of the CMT method makes it possible to supply less heat, which is one of the

factors resulting in a smaller angular deformation and limitation/elimination of hot cracks, if

compared with the MIG-Pulse method applied using the same settings of current-voltage

parameters.

4. Properly conducted heat treatment makes it possible to obtain joints of EN AW 6082 alloy

characterised by strength only slightly lower than that required in case of joints made of steel

grade S355.

References

[1] J. Adamiec, T. Pfeifer, J. Rykała: „Modern methods of aluminum alloys welding”, Solid State

Phenomena, Vol. 176 (2011)

[2] J. Matusiak, T. Pfeifer: Influence of material and technological conditions on quality of welded

joints and emission of fumes into work environment. BIUL.I.S., 2008, no 5.

[3] B. Irving: „Welding the four most popular aluminum alloys”, Welding Journal no. 2/1994

[4] T. Anderson: „Aluminium welding within the automotive industry”, Svetsaren no. 2-3/2001

[5] T. Anderson: „Troubleshooting in aluminium welding”, Svetsaren no. 2/2000

[6] T. Anderson: „How to avoid cracking in Aluminium Alloys”, Welding Journal no. 9/2005

[7] Z. Huda: „Precipitation Strengthening and Age-Hardening in 2017Aluminum Alloy for

Aerospace Application”, European Journal of Scientific Research, no. 4, 2009

[8] N.G. Tretyak, A. Ishchenko, Ya., M.R. Yavorskaya: „Susceptibility of aluminium – lithium

alloys to hot cracking in welding”, Welding in the World no. 1/1995

[9] G. Mrówka-Nowotnik, J. Sieniawski, A. Nowotnik: „ Effect of heat treatment on tensile and

fracture toughness properties of 6082 alloy”, Journal of Achievements in Materials and

manufacturing Engineering, Vol. 32, Issue 2 (February 2009)


Fabrication of ceramic-metal composites with percolation of phases

using GPI

Anna Boczkowska1,a, Paulina Chabera2,b, Anna J. Dolata3,c, Maciej Dyzia4,d, Rafał Kozera5,e and Artur Oziębło6,f

1,2,5 Warsaw University of Technology, ul. Wołoska 141, 02-507 Warsaw, Poland

3,4 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

6 Institute of Ceramics and Construction Materials, ul. Postępu 9, 02-676 Warsaw, Poland

a [email protected],

b [email protected],

c [email protected],

d [email protected],

e [email protected],

f [email protected]

Keywords: porous ceramics, cast aluminium alloy, composite

Abstract Al2O3/AlSi12CuMgNi composites were fabricated using gas-pressure infiltration (T=7000C, p=4 MPa) of an aluminium alloy into alumina performs. Volume fraction of the ceramic phase was up to 30%, while the pore sizes of the ceramic preforms varied from 300 to 1000 µm. Ceramic preforms were formed by method of copying the cellular structure of the polymer matrix. The results of the X-ray tomography proved very good infiltration of the pores by the aluminium alloy. Residual porosity is approximately 1 vol%. Image analysis has been used to evaluate the specific surface fraction of the interphase boundaries (Sv). The presented results of the studies show the effect of the surface fraction of the interphase boundaries of ceramic-metal on the composite compressive strength, hardness and Young’s modulus. The composites microstructure was studied using scanning electron microscopy (SEM). SEM investigations proved that the pores are almost fully filled by the aluminium alloy. The obtained microstructure with percolation of ceramic and metal phases gives the composites high mechanical properties together with the ability to absorb the strain energy. Compression tests for the obtained composites were carried out and Young’s modulus was measured by the application of the DIC (Digital Image Correlation) method. Moreover, Brinell hardness tests were performed. Gas-pressure infiltration (GPI) allowed to fabricate composites with high compressive strength and stiffness.

Introduction

Most of the work on metal matrix composites (MMCs) has been concentrated on the particle or fibre reinforced composites in which the ceramic phase is randomly dispersed or oriented in one or two directions [1-7]. A new class of composite materials, which has been developed recently is termed interpenetrating phase composites (IPCs). Such composites are characterized by two continuous phases, both distributed in three directions. This results in a material with better characteristics, because of the combination of two phases with significantly different properties such as strength and strain, which can be optimized [8-10]. Special attention is given to the ceramics matrix composites, mostly infiltrated by light metal. This class of materials exhibits superior strength, toughness and thermal shock resistance compared to monolithic ceramics. Ceramic-metal composites combine the ceramic high stiffness, high strength at elevated temperature and remarkable hardness and wear resistance with the low elastic modulus, high coefficient of thermal expansion and low wear resistance of the light metals alloys. Such materials can find their application in many industry branches, such as aircraft, automotive and armaments industries, as well as in electrical engineering and electronics. The progress is determined by increasing of exploitation of the engineering materials in special application [11-15].


Ceramic-metal composites with two interpenetrating phases can be obtained by several techniques, in particular by standard pressure or pressure less infiltration of porous preforms, by hot pressing or reactive metal penetration. The important issues concerning the fabrication of ceramic-metal composites are problems of wettability of ceramics by liquid metals and achievement of the strong bonding at the ceramic–metal interface [14-17]. This paper is aims at preparation and characterization of co-continuous ceramic/metal composites obtained using gas-pressure infiltration. In this study, the effect of specific surface fraction of interphase boundaries on mechanical properties of ceramic-metal composites was shown. These composites are currently developed for passive protection elements.

Materials and methods

The ceramic preforms were manufactured in the Institute of Ceramics and Construction Material by sintering of RA-207LS Al2O3 powder supplied by Alcan Chemicals. The chemical composition of aluminum oxide was Al2O3 (99,8wt.%), CaO (0,02wt.%), SiO2 (0,04wt.%), MgO (0,04wt.%), Fe2O3 (0,03wt.%), Na2O (0,07wt.%). For each ceramic preforms the porosity was at the same level, approximately 72vol.%. Porous aluminum oxide preforms were formed by the method of copying the cellular structure of the polymer matrix [18]. Three types of polyurethane sponges, differing in density and size of pores were exploited: 60, 45 and 30 pores per inch (ppi). This results in the fabrication of preforms with pore sizes varying from 300 to 1000µm.

An autoclave to gas-pressure infiltration (GPI) designed and built at the Faculty of Materials Science and Metallurgy, Silesian University of Technology (PL) was applied for infiltration of ceramic preforms by EN AC- AlSi12CuMgNi (AK12) cast aluminium alloy [19-21]. Composition of the alloy used for infiltration was: Si–12 wt.%, Fe–0,44 wt.%, Cu-1,08 wt.%, Mn-0,16 wt.%, Mg-1,28 wt.%, Zn–0,14 wt.%, Ni-1,06 wt.%, Ti-0,03 wt.%, Al – remainder. As a result three kinds of Al2O3/ AlSi12CuMgNi (Al2O3/ AK12 ) composites were obtained (Tab.1).

Table 1. Designation of ceramic-metal composites obtained by gas-pressure infiltration GPI.

Designation of Al2O3/ AK12 composites

Pore size of Al2O3

preforms

Designation of ceramic preforms

Al2O3_1/AK12 300-450µm Al2O3 _1

Al2O3_2/ AK12 400-550µm Al2O3 _2

Al2O3_3/ AK12 800-1000µm Al2O3 _3 The microstructure of fabricated composites was studied using Scanning Electron Microscopy

and quantitatively characterized by image analysis. Such parameters as the volume fraction of phases and the specific surface fraction of the interphase boundaries (Sv) were calculated using the Micrometer software. The microstructure was also characterized using X-ray tomography type SkyScan 1174. Brinell hardness tests were also performed. Furthermore, compression tests were carried out using a Zwick 250 machine with application of Digital Image Correlation (DIC) method. The DIC method was utilized to determine Young’s modulus.


Results and analysis

The volume fraction of the ceramics phase is approximately 28 vol.%, the remaining area (72 vol.%) can be filled up with liquid metal (Tab. 2). The results of X-ray tomography proved very good infiltration of the pores by the metal. The composites obtained from preforms with the smallest pores exhibit the smallest residual porosity (<1 vol.%).

Table 2. Volume fraction of Al2O3 ceramics and pores.

Designation of Al2O3

preforms Volume fraction of Al2O3

ceramics [%] Volume fraction of

pores [%]

Al2O3 _1 30,57 69,43 Al2O3 _2 28,52 71,48 Al2O3 _3 25,58 74,42

Table 3. Degree of infiltration of ceramics by cast aluminium alloy.


Residual porosity [%] Composite volume [%]

Al2O3 _1/AK12 0,57 99,43 Al2O3 _2/AK12 1,60 98,40 Al2O3 _3/AK12 1,44 98,56

Compression tests for the samples of porous ceramics and composites Al2O3/AK12 were carried

out. The character of the stress-strain curves of the ceramic and composites was compared. The values of the compressive strength are shown in Table 4.

Table 4. Surface fraction of the interphase boundaries (Sv), compressive strength, hardness HB,

Young’s modulus and energy absorption of Al2O3/ AK12 composites.


Sv [1/mm]

Pore size of Al2O3

preforms

Compressive strength [MPa]

Hardness HB

Young’ modulus

[GPa]

Energy absorption [MJ/m2]

Al2O3 _1/AK12 10,3 300-450µm 341 97,3±1,2 51,7 15,2 Al2O3 _2/AK12 8,61 400-550µm 317 76,2±2,9 44,1 13,6

Al2O3 _3/AK12 5,96 800-

1000µm 294 72,9±0,8 41,6 11,1

The composites fabricated by infiltration of the preform with the smallest pores are characterized by the highest compressive strength. As shown on Figure 1, composites exhibit much higher compressive strength in comparison to the porous preform. Also the slope of the stress-strain curves for ceramics and composites changed. A distinct decrease in stresses on the stress-strain curves of the composites was not observed, while ceramics failed to pass the test. The obtained microstructure with percolation of the ceramic and metal phases gives the composites high mechanical strength together with the ability to absorb strain energy. Young’s modulus (E) was determined as a slope coefficient of the stress-strain curve within the elastic range. The image analysis have been used to measure the specific surface fraction of the interphase boundaries (Sv) using Micrometer program. The measured Sv parameter increases together with the decreasing of the sizes of ceramics perform pores (Tab. 4). Decreasing of the size of perform pores results in the growth of the volume fraction of ceramics-metal interphase boundaries, at permanent value of the porosity.


0

50

100

150

200

250

300

350

400

0 1 2 3 4 5 6

Strain [%]

Str

es

s [

MP

a]

Al2O3_1/AK12

Al2O3_2/AK12

Al2O3_3/AK12

Preform Al2O3 _1

Fig. 1. Typical stress-strain curves for ceramic and Al2O3/ AK12 composites.

The influence of the specific surface fraction of the interphase boundaries (Sv) on mechanical properties of Al2O3/AK12 composites, such as hardness HB, compressive strength, energy absorption and Young’s modulus is shown in Figure 2.

290

300

310

320

330

340

350

5 6 7 8 9 10

Sv [1/m]

Rc

[M

Pa

]

0

50

100

150

200

250

HB

, E

[G

Pa

],

En

erg

y [

MJ

/m2

]

Compressive strength [MPa]

Young’s modulus [GPa]

Hardness HB 306,5/5

Energy absorption [MJ/m2]

Fig. 2. Effect of the specific surface fraction of interphase boundaries on mechanical properties of Al2O3/ AK12 composites fabricated by gas-pressure infiltration of ceramic performs.

The hardness HB, compressive strength, energy absorption and Young’s modulus increase with

an increase of the Sv parameter. The curves have exponential character. A double increase of the fraction of the interphase boundaries causes closely twice an increase of the mechanical properties of the composites with percolation of the microstructure.

Samples after compression do not lose their cohesion, what is visible in Fig 3. The cracks propagation is observed only in the ceramic phase, they are extinguished or deflected by the metal phase (Fig. 4)


(a) (b)

Fig. 3. Images of Al2O3/ AlSi12CuMgNi composites before (a) and after (b) compression test

Fig. 4. SEM images of microstructure of Al2O3/ AlSi12CuMgNi composites fabricated by gas-

pressure infiltration of ceramic performs Al2O3_3/AK12 The BSE observations of the composites fabricated by gas-pressure infiltration of ceramic

preforms revealed the new phases, which evolved in the whole volume of aluminium alloy. Also it was observed that all pores are fully filled by aluminium alloy (Fig.5). Moreover, precipitated phases differ in shape, size and way of spacing.

Image analysis has been used to evaluate volume fraction of the precipitated phases for each of composites using Micrometer program. Independently of size of ceramics performs pores volume fraction of the precipitated phases is almost the same (15 vol.%).

(a)

(b)

Fig. 5. BSE images of microstructure of Al2O3/ AlSi12CuMgNi composites fabricated by gas-pressure infiltration of ceramic performs Al2O3_1/AK12 (a and b respectively)


The chemical composition of precipitated phases observed in the microstructures of composites were examined using the EDS method (Fig.6, 7). It was found that in all Al2O3/AK12 composites the same phases were precipitated. In the structure of the composites the presence of silicon (Fig. 4b) and phases containing copper, nickel and magnesium has been identified (Fig. 5). The BSE images and result of the EDS studies show that phase of silicon was precipitated on the boundaries and in the aluminium alloy. Its quantity is higher than any other phases.

(a)

(b)

Fig. 6. Chemical composition of phases observed in Al2O3/ AK12 composites: a) microstructure, b) Si phase in point 1.

(a)

(b)

Fig. 7. Chemical composition of phases containing magnesium, copper and nickel observed in Al2O3/AK12 composites: a) chemical composition in point 2, b) chemical composition in point 3.


Fig. 8. Distribution of silicon in Al2O3/AK12 composite.

Fig. 9. Distribution of iron, copper and manganese in Al2O3/AK12 composite.

Distribution of silicon, iron, copper and magnesium is shown in Figures 8 and 9. The precise

evaluation of precipitated phases on the boundaries required preparation of the composite’s sample using focused ion beam (FIB). Next observations by the application of the SE and STEM methods


were performed (Fig. 10). The results of SE and STEM studies confirmed diffusion of silicon from aluminium alloy to substrate of Al2O3 ceramics. It was observed that phases of ceramic and metal created coherent, corrugated interface without voids on the boundaries

Fig. 10. BSE images of microstructure of Al2O3/AK12 composites fabricated by gas-pressure infiltration of ceramic performs (Al2O3/AK12)

Conclusions

Ceramic-metal composites, obtained via pressure infiltration of porous Al2O3 ceramics by cast EN AC- AlSi12CuMgNi (AK12) aluminium alloy are characterized by a large degree of infiltration of pores by the metal. As a result of ceramics infiltration, composites of two interpenetrating phases are obtained. The obtained microstructure gives the composites high mechanical strength together with the ability to absorb the strain energy. The gas-pressure infiltration GPI ensure high mechanical properties and degree of infiltration.

The mechanical properties of the composites depend on the specific surface fraction of the interphase boundaries (Sv) and the degree of infiltration. The composite obtained via infiltration of the ceramics preform with the smallest pores exhibits the highest value of compressive strength, hardness and Young’s modulus. It was found that the energy absorption ability of the composites increases treble with the growth of the fraction of interphase boundaries. Due to combining ceramics and metal, composites with higher mechanical properties compared to porous ceramics can be obtained. Moreover, such composites do not lose their cohesion during compression, while the ceramics samples were totally broken.

It was proved that developed technology of fabrication the composite material with the ceramics matrix infiltrated by aluminium alloy ensures the required microstructure. The pores are almost fully filled by the aluminium alloy.

Acknowledgements

The studies were carried out within the PanCerMet project No. O R00 0056 07: “The passive protection of mobile vehicles (air and land) against the influence of AP bullets” financed by National Center for Research and Development in Poland.


References

[1] Requena G., Degischer H.P., Creep behaviour of unreinforced and short fibre reinforced AlSi12CuMgNi piston alloy, Materials Science and Engineering A 420 (2006), 265–275

[2] Jun D., Yaohui L., Sirong Y., Wenfang L., Effect of heat-treatment on friction and wear properties of Al2O3 and carbon short fibres reinforced AlSi12CuMgNi hybrid composites, Wear 262 (2007), 1289–1295

[3] Kaczmar J.W., Pietrzak K., Włosiński W.: The production and application of metal matrix composite materials, Journal of Material Processing Technology, 106 (2000), 58-67

[4] W. Hufenbach, M. Gude, A. Czulak, J. Śleziona, A. Dolata-Grosz, M. Dyzia “Development of textile-reinforced carbon fibre aluminium composites manufactured with gas pressure infiltration methods” Journal of Achievements in Materials and Manufacturing Engineering, Vol. 35, Issues 2, pp 177-183, (2009)

[5] Śleziona J., Bases of the technology of composites, Publishing company of the Silesian University of Technology, Gliwice 1998, 28.

[6] Sobczak J., Wojciechowski S., Contemporary tendencies of the practical application of metal composites, Composites, 2(2002)3.

[7] Rosso M., Ceramic and metal matrix composites: Routes and properties, Journal of Materials Processing Technology 175 (2006) 364–375

[8] Scherm F., Völkl R., Neubrand A., Bosbach F., Glatzel U., Mechanical characterization of interpenetrating network metal–ceramic composites, Materials Science and Engineering, A 527 (2010), 1260-1265.

[9] Aldrich D.E., Fan Z., Microstructural characterisation of interpenetrating nickel/alumina composites, Materials Characterization 47 (2001), 167– 173

[10] Konopka K., Olszówka–Myalska A., Szafran M., Ceramic–metal composites with an interpenetrating network, Materials Chemistry and Physics, 81 (2003) 329–332.

[11] Poniznik Z., Salit V., Basista M., Gross D., Effective elastic properties of interpenetrating phase composites, Computational Materials Science 44 (2008) 813–820

[12] Chabera P., Boczkowska A., Zych J., Oziębło A., Kurzydłowski K.J., Effect of specific surface fraction of interphase boundaries on mechanical properties of ceramic-metal composites, obtained by pressure infiltration, Kompozyty 11: 3 (2011) 202-207

[13] Pagounis E., Talvitie M., Lindroos V.K., Influence of the metal/ceramic interface on the microstructure and mechanical properties of hiped iron-based composites, Composites Science and Technology, 56 (1996).

[14] Potoczek M., Śliwa R.E., Myalski J., Śleziona J., Metal-ceramic composites obtained by the pressure infiltration of metal into the ceramic preform about the structure of foam, Ores and non-ferrous metals, R54 2009 nr 11.

[15] Szafran M., Konopka K, Rokicki G, Lipiec W., Kurzydłowski K. J., Porous ceramics infiltrated of metals and polymers, Composites 2002, 2, 5, 313.

[16] Binner J., Chang H., Higginson R., Processing of ceramic-metal interpenetrating composites, Journal of the European Ceramic Society, 29 (2009), 837–842.

[17] Chang H., Higginson R., Binner J., Microstructure and property characterisation of 3-3 Al(Mg)/Al2O3 interpenetrating composites produced by a pressureless infiltration technique, J Mater Sci (2010), 45:662–668


[18] Oziębło A., Jaegerman Z., Traczyk S., Dziubak C., Porowata ceramika do wytwarzania kompozytowych materiałów metalowo-ceramicznych metodą infiltracji ciśnieniowej ciekłymi stopami aluminium, Szkło i Ceramika, Rocznik 57 (2006).

[19] A. Dolata-Grosz, M. Dyzia, J. Śleziona: “Manufacture and structure of infiltrated of Al-carbon fibres composites” Archives of Mechanical Technology and Automation Vol. 30, no 3, pp 11-18, 2010.

[20] A. Dolata-Grosz, M. Dyzia, J. Śleziona: Structure of Al-CF composites obtained by infiltration methods, Archives of Foundry Engineering, Vol. 11, Special Issue 2, pp. 23-28, 2011.

[21] A. Dolata-Grosz, M. Dyzia, J. Śleziona: Al/CF composites obtained by infiltration method, Kompozyty (Composites), vol. 4, 2011.


Producing of composite materials with aluminium alloy matrix containing solid lubricants

Andrzej Posmyk 1,a, Jerzy Myalski 1,b

Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland


Keywords: composite material, glassy carbon, infiltration, precursor, pyrolysis, friction coefficient, solid lubricant

Abstract. The paper presents the basics information about manufacturing and selected properties

of composite with aluminum alloy matrix containing glassy carbon as a solid lubricant. The so far

used method based on mixing the prepared glassy carbon particles with a liquid metal matrix, has

been compared with a new method elaborated by the authors of the article. With this novel method

carbon is introduced into a composite with the application of liquid carbon precursor and porous

ceramic foams. It is then followed by precursor pyrolysis where, as the result, glassy carbon is

obtained. Ceramic foams help liquid precursor penetrate the ceramic spheroid pores by forming a

thin film of glassy carbon on their walls. The composite produced in such a way features uniform

distribution of carbon within its entire volume which significantly improves tribological properties

of the composite. Costly mixing procedure is not needed. Sliding friction coefficient of the new

composite against cast iron (µ = 0.06-0.28 at wearing in and 0,12 after wearing in) is much lower

than in case of composite containing only ceramic foam as a reinforcing phase (µ = 0.25-0.32).

Introduction

The modern automotive industry uses many kinds of engineering and lubricating materials. One of

the most important functions of the engineering materials is to reduce the weight of automotive

parts for fuels saving and make the vehicle more dynamic due to mass reduction. New engineering

materials – the composites - have been elaborated and applied for automotive industry since the

‘70s of 20th

century. Composites and hybrid composites with light metal matrices have found their

application in production of engine pistons and cylinder liners and for air compressors supporting

brake systems [1-3]. Composites with polymer matrix have found their application in production of

car body and internal equipment of vehicles.

An important role of lubricating materials is to maintain the vehicle efficiency on high level and

reduce fuel consumption through minimizing the friction forces. Better lubrication means lower

wear of rubbing vehicle parts, which results in longer durability.

Some composite materials allowed the combination of the above mentioned functions of the

materials, i.e. reducing the vehicle mass and lowering the friction losses in vehicle subassemblies

due to the incorporated solid lubricants.

At the Silesian University of Technology two novel composite materials has been developed. The

first one, including glassy carbon particles stochastically distributed in aluminium alloy matrix, has

been produced using conventional mixing of matrix and reinforcing particles. The second one,

hybrid, composite containing aluminum oxide spheroids as a reinforcing phase and deposited on the

spheroid walls glassy carbon as a solid lubricant has been developed using liquid carbon precursor

introduced into the spheroids [4,5,6]. The detailed description of the producing methods can be

found in literature [5,6]. This materials can be used for production of different parts of vehicles for

example piston and cylinder liners of combustion engines and piston air compressors.


Experimental details

Production of composites containing solid lubricants

The eutectic silumine (AC-AlSi12Cu1Ni1Mg) which contains glassy carbon (GC) as a reinforcing

phase has been one of the first examined materials. Particles of prepared glassy carbon with mean

diameter up to 150 µm have been used as reinforcement. In order to ensure proper wettability and

prevent thermal degradation, the surface has been protected with a fail-safe coatings. Those were

nickel coatings deposited on the surface of carbon particles by the method of chemical reduction

from nickel salt solutions as well as sodium hydroxide coatings deposited by soaking glassy carbon

particles in NaOH aqueous solution.

An important function of the coatings was to prevent the reactions between carbon and aluminum

matrix, specially to reduce hydrophilic aluminum carbide which might cause degradation of carbon

in wet environment. Directly before glassy carbon particles are introduced into the matrix, they are

heated in temperature of 300°C for 1 hour to remove water. Prepared in such a way particles are

introduced into liquid alloy while being mixed intensively in argon atmosphere. The obtained

suspension is subjected to gravitational casting into a metal mould. Despite of the fact that the

obtained suspension is homogenous, crystallization processes cause diverse and heterogeneous

distribution of carbon particles in the matrix due to the difference in density of carbon and metal

particles. All this seems to bring a lot of disadvantages for tribological properties of the composite.

Fig. 1 presents an exemplary microstructure of the composite produced with the method. The dark

glassy carbon particles on the matrix alloy are visible. On the left bottom corner is a agglomerate of

glassy carbon particles to see (Fig 1a).

a) b)

Fig. 1. Microstructure of the AC-AlSi12Cu1Ni1Mg+15% glassy carbon coated with Ni composite

produced by mixing method, particle diameter 150 µm

The second tested composite has been produced since it was possible to combine the technology of

gel casting of alumina foams [7] with nanotechnology in production of glassy carbon as well as

high pressure infiltration of aluminum alloys. This technology is used to produce composites with

aluminum matrix reinforced with porous ceramics. The porous foam from aluminum oxide applied

here functions as reinforcing phase but also constitutes a frame – an additional engineering element

which ensures composite stiffness. Furthermore, porous ceramics reduces the composite density and

a) b)


ensures uniform glassy carbon distribution within the entire volume of the composite. Glassy

carbon, which covers the ceramics, functions as a lubricant as in case of AC-

AlSi12Cu1Ni1Mg+15% GC composite, but ensures uniform distribution of carbon in the structure

eliminating failures in casting composites such as formation of particles concentration and

agglomerations. Apart from aluminum alloys, other metal alloys such as magnesium or copper, can

be used as matrix. Composite production consists of the following stages:

1. Producing a ceramic foam with porosity up to 95% which would ensure high wear resistance

and lower composite density, average spheroid diameter 100 µm;

2. Ceramic foam saturation with catalyst which would induce pyrolysis and saturation with

furfuric alcohol as a carbon precursor to ensure low friction resistance and low wear of sliding

pairings;

3. Pyrolysis of carbon precursor in argon atmosphere, (3,5 h by 1000°C in argon atmosphere);

4. High pressure infiltration with aluminum alloy penetrating the foam which has been saturated

with glassy carbon (p = 4 MPa).

The amount of glassy carbon produced on the ceramics walls and its properties depend upon the

production process parameters as well as the materials used. Foam which contains spheroids with

bigger diameters and larger pores enables to introduce more liquid carbon precursor and so does the

longer time of saturation.

The conditions of precursor pyrolysis would decide about such properties of glassy carbon as

hardness and shear strength. In higher temperatures of pyrolysis it is possible to obtain carbon

which features higher hardness and lower shear strength. This again conditions its good lubricating

properties [8].

Pressure and infiltration time of foam with liquid matrix alloy conditions the fact of the foam being

filled with an alloy thus about the density and tightness of the composite. If the filling is not large

enough it might cause brittleness and porosity of the composite. Porous materials are not suitable

for the elements of air compressors because they can causing leakage by higher compressions.

a) b)

Fig. 2. Ceramic foam before (a) and after (b) introducing of glassy carbon (macrophotographs)

Structure of examined materials

It is the structure which decides about the properties of the materials. Therefore the basic

investigations of the structure of composites in individual phases of their production i.e. ceramic

foam before and after saturation with carbon precursor and carbonization as well as after infiltration

with matrix material i.e. AC-AlCuMg1 alloy have been performed. The results are presented in

Figs. 2-5. Fig. 2 presents macrographs of ceramic foam before (Fig. 2a) and after saturation (Fig.

2b) with carbon precursor and its pyrolysis.


a) b)

Fig. 3. Ceramic foam with micropores before introducing of carbon precursor

(visible microporosity of oxide spheroids)

a) b)

Fig. 4. Ceramic foam after pyrolysis of carbon precursor (a) and element analysis on ceramic

surface (b)

Tribological test of examined materials

The examined composites can be used for pistons or cylinder liners in air compressors. Therefore

the tests on their tribological properties have been performed under friction in air conditions. Cubes

with 10 mm sides have been cut from the tested composites (Fig. 6a and 6b) whereas counter-

samples in form of rectangular (14x60x6 mm) have been obtained from cast iron (GJL-300)

cylinder liner of a piston compressor (Fig. 6c). Thanks to tem the friction conditions in sliding

contact between examined composite and cast iron was similar to the really conditions in the

compressor. Fig. 7 presents schematically the friction contact of the laboratory stand which

simulates sliding of a piston skirt against the cylinder liner of an air compressor.


Fig. 5. The hybrid composite with ceramic foam as reinforcing phase and glassy carbon as solid

lubricants and AC-AlCuMg1 –alloy as matrix: a) polished crossection, b) ceramic sphere

wall/glassy carbon border; 1 – matrix alloy, 2 – ceramic reinforcing phase,

3 – glassy carbon layers on ceramic surfaces

Tribological investigations have been carried out under the following conditions: reciprocating

motion with relative velocity of v=2.5 m/s, unit pressure p=2 MPa, sliding time 30 min. Short

sliding time results from several minutes long operation of the compressor to supply the tank with

air, after the braking system was used. During the tests friction coefficient has been measured with

the use of a strain gauge force transducer. The accuracy of the transducer was 3% of measured

value. Table 1 lists the values of friction coefficient. The friction coefficient of composite

containing glassy carbon elaborated by mixing method has been added in the table 1. During

friction in air against cast iron was the friction coefficient high (µ=0.32) therefore the tests was in

limited lubrication conditions (2 mg of Semisynthetic 10W/40 oil) conducted.

For a possibility to determine of the influence of the glassy carbon on the friction coefficient the

comparatively investigations of two contact have been carried out:

- composite containing only matrix and ceramic foam/cast iron,

- composite containing matrix, ceramic foam and glassy carbon/cast iron.

a) b) c)

Fig. 6. Samples for tribological investigations; a) matrix + ceramic foam,

b) matrix + ceramic foam + glassy carbon, c) cast iron countersample


Tab. 1. Dependence between friction coefficient and sliding time of the tested materials

(p=2 MPa, v = 2.5 m/s, friction in air)

Friction coefficient, µ

Sliding time, [min] 0 5 10 15 20 25 30

Composite without GC 0.25 0.30 0.30 0.30 0.32 0.32 0.32

Composite with GC 0.06 0.28 0.18 0.14 0.08 0.12 0.12

Composite with GC produced using

mixing method, lubricated

0.05 0.05 0.05 0.05 0.05 0.05 0.05

Fig. 7. Friction contact of the tester used for tribological investigations: 1- composite sample,

2- cast iron countersample, 3- wear track, v- sliding velocity, F- load

a) SEM b) SEM

c) macrograph d) macrograph

Fig. 8. The surfaces of the tested materials after friction: a) composite without GC, b) composite

with GC, c) cast iron after sliding against composite without GC, d) cast iron after sliding against

composite with GC; visible scratch marks of abrasive wear (1) and fragments of oxide ceramics

included into cast iron (2)


Discussion of results

Figures 2 and 3 show ceramic spheroids with pores from about a dozen to 200 µm diameter which

enable to enter both precursor of glassy carbon and a catalyst in a liquid state. The existing spaces

between spheroids enable their external surfaces to be coated with glassy carbon. Fig. 4 shows the

surfaces of ceramic spheroids coated with thin film of carbon (qualitative analysis, Fig. 4b - small

C-pick). The carried out investigations proved the possibility of carbon precursor to be introduced

and its pyrolysis upon the surfaces of ceramic spheroids to be performed. As the result thin carbon

films develop (Fig. 5b). Inside volume of most of spheroids with diameter of 100 up to 500 µm has

been filled with matrix alloy (white fields in Fig. 5a). Inside of some spheroids a bundle of glassy

carbon have been deposited (black area in Fig. 8b). The presence of argon at pyrolysis does not

cause any side reactions upon the carbon surface which could worsen its tribological properties. At

the process of sliding, carbon wear debris are deposited on the surface of cast iron matrix thus

reducing the friction .

However, the produced composites feature some defects like low porosity which has been detected

at macroscopic and macroscopic (Fig. 6) examinations of the cubes surfaces designed for

tribological tests. Both the composite with ceramic foam (Fig. 6a) and ceramic foam plus glassy

carbon (Fig. 6b) demonstrate only slight microporosity. The porosity is visible as small black

cavities on the surface after polishing. Cavities in composite including glassy carbon are bigger then

in composite with ceramic spheroids only. This is caused by lower wettability of carbon surface by

liquid matrix. These cavities can be used as oil deposits.

In order to eliminate this porosity it is necessary to undertake optimization investigations of the

manufacturing process, first of all of the high pressure infiltration with matrix aluminium alloy.

Scanning microscope studies of composites after infiltration with aluminum alloy proved the

presence of thin glassy carbon films upon ceramic spheres walls. Their thickness is varied from 5 to

15 micrometers. The thickness of carbon layer depends among others on the amount and diameter

of pores in the spheroids as well as the duration of saturation with a carbon precursor. Spheroids

with smaller pores include les precursor and the carbon layer is thinner.

Tribological investigations of the produced composites sliding against cast iron (GJL-300) showed

that the presence of glassy carbon films on the spheroids walls reduces friction coefficient

(minimum value µ=0.06). The time of initial wearing-in is 20 min. In the initial phase of sliding

small parts of oxide ceramic spheroids crash and crack (Fig. 8b) and enter into the sliding surface of

cast iron (Fig. 8d). This is accompanied by higher friction coefficient – constant during sliding of

composite without glassy carbon and temporary for 10 min in case of sliding of composite with

glassy carbon. Fig. 8b shows on the background of matrix material the revealed inside of the oxide

spheroid filled with glassy carbon bundles. This carbon has not yet took part in friction because

there are no traces (scratches) of sliding visible. Petty cracks of matrix material which result from

deformation of aluminum oxide walls can be spotted in the vicinity of crashed spheroid. Since such

spheres filled with carbon are distributed upon the entire friction surface and at varied distances

from this surface, they function as solid lubricant depots. The depots open as the wear of matrix

material and ceramics walls is progressing.

In case of comparative contact, e.g. without glassy carbon, friction coefficient stabilizes after about

10 minutes and is 0.3. Here, as in above mentioned contact, some parts of oxide crumble and get

included into cast iron rubbing surface (Fig. 8c). Fig. 8a presents composite surface without GC

after friction. There are cracks on the matrix surface which result from pressure and deformation of

oxide spheroids walls. In consequence crushing of oxide walls and inclusion of fragments in cast

iron occurred (Fig. 8c). These fragments cause abrasion wear of matrix material. Friction forces in

this contact are large enough to generate slight plastic deformation of the matrix material visible in

Fig. 8a in the form of gentle, parallel lines. A local plasticizing of matrix material is caused by heat

concentration resulting from lower heat conductivity of oxide ceramic placed under the matrix

surface (~20 W/mK for Al2O3 and ~150 W/mK for matrix alloy).


Conclusion

As the result of the carried out investigations it has been found that the production of a hybrid

composite by introducing carbon precursor into porous ceramic preforms and the pyrolysis which

follows is possible. An improvement of tribological properties of so far manufactured composite

with aluminum cast alloy matrix (AC-AlCuMg1) containing foamy oxide ceramics has been

achieved. Coefficient of friction in air sliding against cast iron GJL-300 composite which contains

only ceramics is 0.3, whereas composite after glassy carbon is added equals 0.12. When carbon is

placed inside ceramic spheroids its uniform distribution over the entire volume is possible. This

prevents local tacking between composite matrix and a sliding partner in places where less number

of glassy carbon particles observed which happened during sliding of composite produced by

mixing method.

The composite produced by introducing of carbon precursor into ceramic spheroids show better

tribological properties as the composite produced heretofore by mixing of matrix alloy with

prepared glassy carbon particles because of better distribution of glassy carbon on the walls and

inside of the oxide ceramic spheroids.

High pressure infiltration process of the shaped ceramic foam containing glassy carbon with liquid

matrix alloy requires still more optimization in order to decrease ore eliminate local porosity.

References

[1] M. Dyzia, AlSi7Mg/SiC and Heterophase SiCp+Cg Composite for Use in Cylinder-Piston

System of Air Compressor, Solid State Phenomena Volume 176 (2011) 49-54.

[2] A. Dolata-Grosz, Interaction of Al-Si alloys with SiC/C ceramic particles and their influence on

microstructure of composites, Solid State Phenomena Volume 176 (2011) 55-62.

[3] M. Dyzia, A. Dolata-Grosz, J. Wieczorek, J. Sleziona: Die-cast heterophase composites with

AlSi13Mg1CuNi matrix, Archives of Foundry Engineering, Vol. 10, 1/2010 301-304

[4] A. Posmyk, Myalski J., Wistuba H., Producing of composite materials containing solid

lubricants. Polish patent application [WIPO ST 10/C PL 398311].

[5] Myalski J., Formation of tribological properties of glassy carbon composites. Publisher Silesian

University of Technology, Gliwice 2011.

[6] Producing method of tribological composite material with aluminium matrix. Polish Patent No

197636.

[7] M. Potoczek, J. Myalski, J. Sleziona, R.E. Sliwa, Gelcasting of alumina foams as preforms for

metal infiltration. Material Engineering No 6 (172), 2009, 536-539.

[8] W.V. Kotlensky, D.B. Fischbach, Tensile and structural properties of glassy carbon. NASA

Technical Report No 32-842.


Machinability of aluminium matrix composites

J. Wieczorek1, a, M. Dyzia1,b and A. J. Dolata1,c 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland.


Keywords: metal matrix composites, aluminium alloy, machinability, surface geometry.

Abstract. The today's interest in MMCp results from a number of their creative properties, which can be designed through a proper selection of reinforcing components and technological parameters. The composite machine elements such as engine, compressor parts obtained by casting methods require the specially final machining. The introduction of hard ceramic particles increase the wear resistance of composite material compared to non reinforcement alloy. Simultaneously increase wear and reducing the durability of tools cutting. The presence of ceramic particles (SiC, Cs) in aluminium matrix influence on surface geometry formed in track of processing. In this paper the results of investigations of geometry surface of composite after machining. Applied machining conditions for composite material were the same as for unreinforcement alloy, it made possible to compare the conditions of machining processing. It the piston skirt was conducted light profilometry investigation were the parameters 2D and 3D surface topography evaluated. Results shows dependency of surface parameters (Ra, Rz) after machining on kind, size and volume fraction of reinforcement particles applied in composite material.

Introduction

Possibility of producing machine parts and sub-assemblies from MMC composite materials depends on solving different technical and technological problems. These problems need using modern solutions starting from the production process of the material, through its processing up to the final shaping of the ready material [1]. In case of the composites with aluminium alloy matrix reinforced with a ceramic phase in the form of particles there is a necessity of working out the production technology. One of the methods to produce the MMCp composites usable on the industrial scale is mechanical mixing [2]. Created composite suspension is converted into semi-products usually by casting methods. Regardless the type of casting: gravity, pressure moulding, pressure, the composite casts must be subjected to the finishing mechanical treatment. The scope of this treatment depends on the casting technology and includes the following tasks: cutting-off the riser heads and the pouring system , preliminary machining, finishing treatment. In each stage of the composite cast treatment the problem of machining arises. Machining defined as the material susceptibility to form in scope of the machining includes: cutting, turning, drilling, milling, grinding, polishing. From the industrial point of view, the notion of machining ability also includes durability of machining tools, energy costs and time needed to perform all the necessary technological operations [3]. In case of machining the composites reinforced with ceramic particles, the main issue is the type and morphology of used reinforcement which limits the machining ability. This problem appears because these composites were designed regarding the highest resistance to wear in friction conditions so these features are in opposition to machining. Nowadays metal-ceramic composite materials are used to manufacture car parts such as brake discs, pistons, brake shoes, belt pulleys [4]. The greatest disadvantage of MMC composites is the high cost of finishing machining. The basic problem that appear are intensive wear of the tool point, deformation, cracking and other damages of the reinforcement phase of the machined material. These kind of damages may appear due to plastic deformation of the material machined with the tool wear. The aim of the composite casts machining is to give a ready-made product proper shape and size as well as forming the proper geometrical surface [5]. The experiments carried out in the industrial conditions including cutting-off riser process from the composite pistons followed by machining through turning revealed that


the considerable technical problem is caused by the low machining ability of the composites. The phenomenon of intensive wear of machining tools and difficulty in proper surface forming of the final MMC composite product caused by it became the reason for starting the research whose results are presented in this article.


Wrong choice of machining parameters may lead to damage of the cast surface and machining tool. The sample of this effect is shown in figure 1 that presents a part of composite roll surface machined with sintered carbides. The surface is wrongly machined, one can notice the damages and deep scratches on the surface. Figure 1b presents the state of the tool after machining. The tool is totally damaged due to ceramic phase reinforcement of composite.

Fig. 1. Composite roll after machining with sintered carbide tool (a). The machining blade damaged due to influence of composite reinforcing particles (b).

The main reason for the experimental research was the damage of the machining tools and bad quality of the composite cast surface treated with machining. The aims of this research ware to find the proper machining tool, which will give good machining results as well as defining the influence of the composite reinforcing phase on the surface quality after machining process. The basic assumption of the carried research was the assessment of the geometrical features of the composite surface after turning conducted with the parameters similar to those used in the industrial environment at the machining of aluminium alloy products. The subject of the research were composites with aluminium alloy cast matrix AlSi7Mg reinforced with silicon carbide particles SiC with the 25 µm diameter and the mixture of SiC reinforcing particles and glassy carbon. The phase composition of the examined composites is shown in Table 1.

Table 1. Identification of samples.

Samples Matrix Reinforcing

particles Volume fraction of

reinforcing particles, % Reinforcing particles

dimension, µm AlSi7Mg AlSi7Mg - - -

AlSi7Mg+SiC AlSi7Mg SiC 15 25

AlSi7Mg+SiC+Cs AlSi7Mg SiC;

Cs (glassy garbon) 10 5

25 50

Research on machining ability of the composites were conducted on the basis of the gradual trial of turning made with the rolls casted from the composite. The mechanical conditions of the conducted experiment are as the following: turning velocity S=1400 turns/min and the move f=0,14. The PCD (polycrystalline diamond) blade was used for machining. The initial 25mm diameter rolls were turned up to 1mm diameter- the turning considered 11 diameters, at each turn the diameter was reduced by 14 mm (Figure 2).

a) b)


Fig. 2. The composite roll after gradual machining trial

The roll surfaces were profilegraphometricaly examined after machining. The aim of this assessment was quality of the task and of the gained roughness. The surface of the machined rolls was also microscopically examined. These examinations were conducted for each of 11 machining zone for each roll. The comparison of the roughness layout in the particular machining spheres show significant differences appearing on the surface of the material while machining depending on the type of used reinforcement. In Figure 3 the roughness profile from the 6th zone is shown for each of the examined material.

0 1 2 3 4 5 mm

µm

0

10

20

30

0 1 2 3 4 5 mm

µm

0

10

20

30

0 1 2 3 4 5 mm

µm

0

10

20

30

Fig. 3. Roughness profile in the 6th zone of the examined rolls: a) AlSi7Mg; b) AlSi7Mg+SiC;

c) AlSi7Mg+SiC+Cs The greatest roughness appears on examined surface of the heterophase composite reinforced with the mixture of silicon carbide and glass carbon particles. The maximal value of the roughness was Ra = 24,5 µm. A bit lower value of roughness 20 µm appears in the composite reinforced with silicon carbide particles (Fig. 3b). The lowest roughness with the maximal value of 10 µm, approximately three times lower than in case of heterophase composite, was measured on the surface of Al7SiMg alloy (Fig. 3a).

a)

b)

c)


Moreover, the uneven layout of the distances between the peaks of roughness in case of the machined composite surface contrary to the even layout measured on the non-reinforced alloy (Fig. 3a). Comparison of the roughness parameters of the casted rolls from the various composite materials with the comparison with the non-reinforced matrix alloy is shown on the Figure 4.

1,082,26 3,22

6,53

13,2

18,4

7,92

18

24,8

0

5

10

15

20

25

AlSi7Mg AlSi7Mg+SiC AlSi7Mg+SiC+Cs

µµ µµm

Ra Rz Rt

Fig. 4. Comparison of the roughness parameters (Ra, Rz, Rt) of the composite materials after the

machining trial. 3D picture layout of roughness and microscopic photos of surfaces after machining shown in Figures 5 - 7 allow to recognize caused of differences among measured geometrical parameters on the surface after machining. In case of the matrix material, the work of the blade is not interrupted – trace of machining is regular and the roughness parameters are proportional to the set machining parameters. The distances between the roughness peaks are set by the quotient of the feed rate (0,14µm) and spindle cutting speed (1400 rpm) (Figure 5).

Fig. 5. Geometry surface of AlSi7Mg alloy after machining: a) 3D image of the surface, b) microscopic image

On the surface of the SiC particle reinforced composite irregular scratches appear, their deepness is 20 µm. Hard particles of the ceramic reinforcement are the obstacle in the work of the blade, they cause its vibrations which are the reason for the uneven roughness profile on the machined surface. Some of the reinforced particles are torn away from the matrix during the machining. This the additional element having influence on shaping the surface. The trace after tearing the particle away is visible in the microscopic photo (Figure 6).

a) b)


Fig. 6. Geometry surface of AlSi7Mg + SiC composite after machining: a) 3D image of the surface,

b) microscopic image Using the particle mixture of SiC and glass carbon causes another change of the machining ability. Brittle glass carbon particles become cracked in contact with the machining blade. The crushed particle of glass carbon is removed from the matrix and the remaining crater can be observed on the machined surface (Fig. 7a). The powder that results from the damaged particles of glass carbon gets between the machining blade and the machined material becoming a lubricant. Its presence causes sliding of the machining tool observed during the performance of machining.

Fig. 7. Geometry surface of AlSi7Mg + SiC+Cs composite after machining: a) 3D image of the surface, b) microscopic image

In comparison with the new blade the machining edge after turning was unspoiled. The obtained results are shown in Figure 8. The only visible trace of the cooperation is the tiny amount of the matrix material that adhesively stays on the machining edge (Figure 8b). This effect is typically observed on the surface of the machining tools after the performance with the matrix aluminium alloys materials.

a)

a)

b)

b)


0 0.1 0.2 0.3 0.4 0.5 mm

µm

0

20

40

60

80

0 0.1 0.2 0.3 0.4 0.5 mm

µm

0

20

40

60

80

Fig. 8. Comparison of the diamond machining blade surfaces before a) and after turning b).

Summary

Introducing of the aluminium alloy reinforcement in the form of ceramic particles significantly influences the machining ability of the composite material. Machining of the composite materials demands the usage of the highest quality tools. The satisfying results were only gained when the diamond blade was used. The type of used reinforcement conditions the surface geometry after turning. If we use the silicon carbide as the reinforcement, we have to assume that the scratches will appear on the surface due to the tearing away of the reinforcement particles. A hard ceramic phase of the reinforcement can also cause crushing of the machining blade. This fact increase the cost of the machining. When we use the mixture of silicon carbide and glass carbon particles as the reinforcement, the greatest influence on the surface forming during machining has the glass carbon. The carbon material crushes and tears away during the machining leaving the visible traces on the machined surface. During the process of glassy carbon particle crushing the obtained usage products reduce the possibility of damaging the machining blade but at the same time causing its slide reduce its efficiency. The results obtained in the turning trial measured usage of machining tools and the surface quality after turning point the possibility of the optimal selection of the machining parameters depending on the type of the material. This needs yet separate detailed research.

References

[1] M. Dyzia, A. Dolata-Grosz, J. Wieczorek, J. Śleziona: Die-cast heterophase composites with AlSi13Mg1CuNi matrix, Archives of Foundry Engineering, Vol. 10, Special Issue 1/2010, p. 301-304.

[2] A. Dolata-Grosz: Heterophase composites, production, properties and structure, Foundry Research Institute, Kraków, 2009, p. 33.

[3] J. Stósa, Machining - innovation, The Institute of Advanced Manufacturing Technology, Kraków 2008.

[4] W. Olszak, The geometric structure of the surface. Machning, WNT, Warszawa 2008, p. 136-147.

[5] A. Posmyk, H. Wistuba, P. Falkowski: Ceramic-Carbon Composites designed for piston group of combustion engines. Composites. 11, vol. 2/2011, p. 97-101.


Influence of particles type and shape on the corrosion resistance of

aluminium hybrid composites

Anna J. Dolata1,a, Maciej Dyzia1,b, Witold Walke2,c 1Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

2 Silesian University of Technology, ul. Akademicka 2A, 44-100 Gliwice, Poland

a [email protected],


c [email protected]

Keywords: aluminium composites, silicon carbide, glassy carbon, hybrid system, corrosion resistance

Abstract AMCs due to good thermal and tribological properties, they are applied as the material for: pistons in modern combustion engines, drive shafts, shock absorber cylinders and brake nodes. Heavy-duty operation, especially under tribological conditions, frequently in corrosive environment, requires knowledge on their corrosion resistance. This paper presents the initial results of the research on susceptibility of aluminium alloy matrix composite material reinforced by SiC particles and mixture of SiC+C particles to corrosion. The purpose of the research was to determine the influence of reinforcing phases, their type and shape on corrosion behaviour in a typical corrosion environment, with low NaCl concentration, in relation to the matrix alloy. Determination of corrosion resistance of Al/SiC+C hybrid composite is a new issue and falls within the field of interest of the authors of this article.

Introduction

Composite materials are more and more often used as construction materials, particularly in the automotive and engineering industry. Wide prospect of application opens before light metal matrix composites with high-strength reinforcement [1]. The important position in this group is held by aluminium alloy matrix composites with SiC, Al2O3 and carbon or graphite particles as well as new material solutions – Al/SiC+C and Al/Al2O3+C hybrid composites [2-4]. The investigations carried out for this group of materials have confirmed that, in comparison with matrix alloy, they have significantly higher stability and rigidity, both at room and elevated temperature, increased wear resistance, better fatigue characteristics, substantially reduced thermal conductivity and reduced coefficient of thermal expansion [5]. Extra properties, such as stabilisation of the coefficient of friction, and above all significant reduction in wear of material coupled with composite in the friction pair, are demonstrated by hybrid systems. Due to their properties, they are applied as the material for: pistons in modern combustion engines, drive shafts, shock absorber cylinders and brake nodes [3]. Heavy-duty operation, especially under tribological conditions, frequently in corrosive environment, requires knowledge on their corrosion resistance [6]. Literature published in this area is limited and often contradictory [4-10]. The problem related to corrosion resistance of Al/SiC+C hybrid composites is a new subject and has never been taken up before. As it results from literature, composite materials may show lower corrosion resistance as compared to aluminium alloy they are made of. It was found that reduction in corrosion resistance may be affected by the following, but not limited to: 1. porosity, dislocation density, as well as stresses at the reinforcing phase/matrix interface [8,9]; 2. chemical and phase composition of both matrix alloy and interface surface; 3. susceptibility to the formation of galvanic pairs between active aluminium alloy and more noble reinforcement material [10,11],


4. existence of intermetallic phase precipitations in composite structure, which is, as demonstrated, favourable for the formation of well conducting “cathode spots” on the passive layer [12]. From the point of view of corrosion resistance, the way of manufacturing composite materials and the shape and size of reinforcing phase – fibre, particle, sphere, are also important [13]. The discrepancies and effects may be a result of different chemical composition of matrix, as well as properties of reinforcing phases. This paper presents the initial results of the research on susceptibility of aluminium alloy matrix composite material reinforced by SiC particles and mixture of SiC+C particles to corrosion. The purpose of the research was to determine the influence of reinforcing phases, their type and shape on corrosion behaviour in a typical corrosion environment, with low NaCl concentration, in relation to the matrix alloy. Determination of corrosion resistance of Al/SiC+C hybrid composite is a new issue and falls within the field of interest of the authors of this article.

Methodology of the research

The tests were carried out for AlSi7Mg aluminium alloy and composite samples: AlSi7Mg/SiC and AlSi7Mg/SiC+C. The structure of aluminium matrix alloy and composites in cast state was characterized by means of light and scanning microscopy methods. Selected, representative structures shown in Figures 1 and 2. For testing corrosion purposes, samples with diameter of d = 13 mm and surface roughness of Ra = 0.50 µm, which was obtained by mechanical working (grinding with water abrasive paper, granularity 500 grains/mm²), were used (Fig. 3.).

a b

Fig. 1. Structure of AlSi7Mg aluminium matrix alloy: a) mag. 100 x, OM, b) mag. 500x.

a b Fig. 2. Structure of composites: a) AlSi7Mg/SiC; b) AlSi7Mg/SiC+C.


Fig. 3. Samples used for testing corrosion obtained by mechanical working.

The pitting corrosion resistance was evaluated based on recording of anodic polarization curves

by potentiodynamic method, using the Radiometr’s electrochemical testing system VoltaLab® PGP 201, which was a part of the measuring system. As the reference electrode, the saturated calomel electrode (SCE) KP-113 was used. The auxiliary electrode was platinum electrode PtP−201. It is one of the basic methods for determination of corrosion resistance of metallic materials [14]. The research included the evaluation of samples for resistance to pitting corrosion. Before the tests, every sample was cleaned in 96% ethyl alcohol in an ultrasonic cleaner. The tests started with determination of opening potential EOCP. Then, the anodic polarization curves were recorded, starting the measurements from potential Estart = EOCP − 100 mV. The change in potential took place in anodic direction with the rate of 1 mV/s. Upon obtaining the plate current density of 1mA/cm², the direction of polarization was changed. In this way, the return curve was recorded. The tests were carried out in 0.01M of NaCl at solution temperature of T = 21±1 °C. From the recorded curves, the characteristic quantities describing resistance to pitting corrosion were determined, i.e.: corrosion potential Ecorr (mV), polarization resistance Rp (Ωcm²), and corrosion current density icorr (A/cm²). For determination of polarization resistance Rp, the Stern method was used. Polarization resistance depends above all on corrosion current, therefore it was assumed in the paper that values β for cathodic and anodic reactions are the same and equal to 0.12 V. The corrosion current density was calculated from simplified dependence icorr = 0.026/Rp.

To obtain information on physicochemical properties of sample surfaces, tests using the electrochemical impedance spectroscopy were carried out [15]. The measurements were taken using the measuring system Auto Lab PGSTAT 302N equipped with FRA2 (Frequency Response Analyser) module. This measuring system allowed testing to be conducted within the frequency range of 104 ÷ 10-2 Hz. The sinusoidal voltage amplitude of the activating signal was 10 mV. The impedance spectra of the system were determined and the obtained measurement data were adjusted to the equivalent system. The impedance spectra of the tested system were presented in the form of Nyquist diagrams for different frequencies and in the form of Bode diagrams. The obtained EIS spectra were interpreted after their adjustment to the equivalent electrical system by the least squares method.


The anodic polarization curves determined for samples with ground surfaces (Ra = 0.50 µm) are presented in Figure 4. It was found from the measurements that average value of corrosion potential for AlSi7Mg matrix samples was Ecorr = -563 mV. The 7% content of silicon carbide (SiC) particles in aluminium alloy resulted in reduction in average value of corrosion potential to Ecorr = -626 mV. In turn, for samples of hybrid composites with 7% content of SiC and 3% content of glassy carbon particles, corrosion potential was Ecorr = -695 mV. In addition, the polarization resistance Rp and corrosion current density icorr were determined, by Stern method, for individual variants of tested samples and they amounted to, respectively:


• AlSi7Mg (AK7) samples – Rp = 49.07 kΩcm2, icorr = 0.53 µA/cm2, • AlSi7Mg /SiC (AK7+SiC) samples – Rp = 34.72 kΩcm2, icorr = 0.75 µA/cm2, • AlSi7Mg /SiC+C (AK7+SiC+S1) samples – Rp = 37.42 kΩcm2, icorr = 0.69 µA/cm2.

Fig. 4. The anodic polarization curves for investigated samples.

The impedance spectra recorded for every tested sample with ground surface are presented in Figures 5-7. The impedance of the electrode/surface layer/solution phase interface was characterized by approximation of the experimental data using the electric model of the equivalent circuit in Figure 8. It has been found that the best adjustment of the experimental impedance spectrum with software-generated model curve for the real and imaginary component of circuit impedance according to changes in the measuring signal for AlSi7Mg matrix alloy sample is obtained by using the equivalent electric circuit comprised of the parallel system including Cp capacity element connected with Rp resistance, characterizing the surface layer, and with resistance at high frequencies, which may be attributed to electrolyte resistance Rs (Fig. 8a and Tab. 1). Later on, the composite samples were tested. The addition of SiC particles and hybrid mixture of SiC+C particles affected the change in layer nature. It has been found that the best adjustment of the experimental impedance spectra (Fig. 8b) is obtained by using the equivalent electric circuit comprised of the parallel system including capacity element connected with resistance of transition and additional element, the so-called Warburg impedance (W), which reproduces the effect of reagents on corrosion, and with resistance at high frequencies, which may be attributed to the electrolyte ohm resistance (Tab. 1).


Fig. 5. Impedance spectrums for AlSi7Mg aluminium matrix: a) nyquist diagram, b) bode diagram.

Fig. 6. Impedance spectrums for AlSi7Mg/SiC composite: a) nyquist diagram, b) bode diagram.

Fig. 7. Impedance spectrums for AlSi7Mg /SiC+C hybrid composite: a) nyquist diagram, b) bode diagram


a) b)

Fig. 8. Physical model for: a) AlSi7Mg – oxide layer – NaCl,

b) AlSi7Mg +SiC – oxide layer – NaCl.

Table 1. Results of EIS

Material Rs,

Ωcm2 Rp,

kΩcm2 Cp, µF

W, Ω-1cm−2s−n

AlSi7Mg 25 13.20 4.33 - AlSi7Mg/SiC 26 23.45 7.29 0.4099e-3 AlSi7Mg/SiC+C 25 14.66 11.69 0.2200e-3

Summary

The performed potentiodynamic tests revealed slight differences in corrosion resistance of tested materials. The value of polarisation resistance Rp was observed to run at similar level, regardless of what type of reinforcement was used. It is caused by the presence of aluminium, i.e. an element with high chemical activity. It shows a tendency to passivation, which provides its alloys with high corrosion resistance in low-aggressive environments. The effective protective barrier in the form of an oxide layer with thickness of only 1 nm is then formed on the surface of aluminium alloy products. Occurred as a result of self-passivation in 0.01 M of NaCl solution, the primary passive layer caused differences in corrosion potential values. AlSi7Mg alloy was found to have higher tendency to self-passivation as compared to composite materials. The electrochemical impedance spectroscopy tests were carried out as a supplement. Sectors of semicircles starting in the origin of the coordinate system are visible on all the Nyquist diagrams (Fig. 5a, 6a and 7a). They indicate the activation control of corrosion processes of tested materials. More deformed fragments of semicircles in the range of low frequencies indicate a significant influence of mass transport on corrosion processes. The effectiveness of corrosion protection first of all depends on the passive layer thickness and tightness. The comparable impedance modules determined for the tested materials also prove their comparable protective properties. The results of impedance spectroscopy experiments on the tested corrosion systems are also presented by means of Bode diagrams (Fig. 5b, 6b and 7b). In all the tested corrosion systems, the phase shift angle was changing as frequency was changed. For AlSi7Mg alloy and AlSi7Mg/SiC composite the distinct maximum occurs in the range 540 ÷ 560 for low frequencies (60 ÷ 90 Hz), while for AlSi7Mg /SiC+C hybrid composite the characteristics reveals the maximum phase angle value equal to 450 for low frequency – 40 Hz. The performed tests revealed that in the environment of 0.01 M NaCl solution the particles of the reinforcing phase in AlSi7Mg alloy matrix showed different values of standard potential, with regard both to aluminium itself and to each other, which results in their different electrochemical reactions and, in the end, different corrosion behaviors. In addition, the mutual reactions between potentials of the composite structural elements have significant influence on the corrosion effect. Local voltaic cells may occur between the reinforcement and the matrix. To provide full description of corrosion properties of the tested materials, further investigations in the environment similar to the real working conditions and microstructure investigations are planned.


Acknowledgements

Scientific work financed from funds allocated for National Science Center, Project No. N N508 630 540

References

[1] D.B. Miracle: Metal matrix composites – From science to technological significance, Composites Science and Technology 65 (2005) 2526–2540.

[2] J. Wieczorek, J. Śleziona, A. Dolata-Grosz, J. Myalski, Influence of material of friction partner on tribological properties of heterophase composites, Kompozyty (Composites) 8:1 (2008) 5-10.

[3] M. Dyzia, AlSi7Mg/SiC and Heterophase SiCp+Cg Composite for Use in Cylinder-Piston System of Air Compressor, Solid State Phenomena Vol. 176 (2011) 49-54.

[4] A. Dolata-Grosz, Interaction of Al-Si alloys with SiC/C ceramic particles and their influence on microstructure of composites, Solid State Phenomena, Vol. 176 (2011) 55-62.

[5] A. Posmyk, J. Cybo, Abrasive and frictional wear properties of reinforced aluminium alloys, Tribologie und Schmierungstechnik, Vol. 44, Issue 2 (1997) 79-83.

[6] A. Posmyk A., R. Czech, The corrosion influence on using of metal matrix composites brake discs, Tribologia, Vol. 43 No 1 (2012).

[7] J. Bieniaś, Aluminium matrix composite materials – Chojen structure and corrosive aspects, Materials Engineering 3 (2006) 561-654 (in Polish).

[8] J.M.G. De Salazar, A. Ureña, S. Manzanedo, M.I. Barrena, Corrosion behaviour of AA6061 and AA7005 reinforced with Al2O3 particles in aerated 3.5% chloride solutions: potentiodynamic measurements and microstructure evaluation, Corrosion Science 41 (1998) 3 529-545.

[9] H.J. Greene, F. Mansfeld, Corrosion Protection of Aluminium Metal-Matrix Composites, Corrosion 53 (1997) 12 920-927.

[10] L.A. Dobrzański, A. Włodarczyk, M. Adamiak, Structure, properties and corrosion resistance of PM composite materials based on EN AW-2124 aluminum alloy reinforced with the Al2O3 ceramic particles, Proceedings of the 13th International Scientific Conference “Achievements in Mechanical and Materials Engineering” AMME’2005, Gliwice-Wisła (2005) 203-206.

[11] L.H. Hihara, R.M. Latanision: Galvanic Corrosion of Aluminium-Matrix Composites, Corrosion 48 (1991) 7 546-552.

[12] M. A. Malik, H. Bala: Electrochemical methods of stability evaluation of passive layers on Al-based composite materials, Material Engineering 5 (1995) 133-137 (in Polish).

[13] K. Łuczak, P. Liberski, J. Śleziona: Influence of volume fraction and particles size on the corrosion resistance of aluminium composite with ceramic particles, Kompozyty (Composites) 3(2003) 6 75-79.

[14] J. Przondziono, W. Walke, J. Szala, E. Hadasik, J. Wieczorek: Evaluation of corrosion resistance of casting magnesium alloy AZ31 in NaCl solutions. IOP Conference Series: Materials Science and Engineering 22 (2011), art. no. 012017.

[15] M. Kaczmarek, W. Walke, Z. Paszenda: Application of electrochemical impedance spectroscopy in evaluation of corrosion resistance of Ni-Ti alloy, Electrical Review 12b (2011) 74-77.


Course of solidification process of AlMMC – comparison of computer

simulations and experimental casting

Roman Zagórski1,a, Anna J. Dolata1,b,Maciej Dyzia1,c 1Silesian University of Technology, ul. Krasińskiego 8, 40-019, Katowice, Poland

a [email protected] ,


c [email protected]

Keywords: alloy casting, solidification, contact resistance, CFD simulations

Abstract. The aim of the paper is to present the possibilities of computational simulations for the casting of aluminum matrix composite (AlMMC) reinforced with ceramics based on experimental data. The comparison of simulation and experimental results concerned the solidification process i.e. the course of solidification, temperature distribution and final arrangement of reinforcement particles. First, we have performed the experimental gravity casting of the aluminum matrix alloy AK12 (AlSi12CuNiMg2) and the composites AK12/SiC and AK12/Cg reinforced with silicon carbide SiC and glass carbon Cg, respectively, into the sand mold. During the experiment we have recorded the temperature using the ThermaCAMTME25 photometer system as well as in the selected point inside the sand mold. Using experimental data we have carried out the numerical calculations according to the methods and procedures contained in the program ANSYS Fluent 13. We have based the simulations on the two-dimensional model in which the Volume of Fluid (VOF) and enthalpy methods have been applied. The former is to describe two-phase system (air-composite matrix free surface, volume fraction of particular continuous phase) and the latter shows modeling of the solidification process of the alloy and composite matrix. We have used the Discrete Phase Model (DPM) to depict the presence of reinforcement particles. The assumption of the appropriate values of simulation parameters has shown that the simulation results are convergent with experimental ones. We have observed a similar course of the composite solidification (temperature change at the designated point), the temperature distribution and the arrangement of reinforcement particles for the simulation and experiment.

Introduction

The theoretical investigations and computer simulation currently play an important role in the researches on composite casting processes. Another important element is to compare the simulation with the experiment. Conducting of computer simulations is closely related to the possession of the relevant experimental data as well as selecting of appropriate tools, models, theories, methods and calculation parameters. Presented computational and experimental results concern the casting of the metal matrix composite reinforced with ceramic particles. This type of material is often the object of theoretical and experimental researches [1-8].

Currently, there are several programs which belong to Computational Fluid Dynamic group (CFD). Most of them apply the methodology of discretization of the area in which the calculations are performed and application of algorithms for solving equations of fluid and heat transfer and solidification to describe the phenomena occurring during the casting process. Fluent, which is one of them, belongs to unified working environment ANSYS Workbench 13 [7-9].

Fluent is based on the assumption of the final volume method. In the elaborated model we have used several methods and techniques contained in this program. The modeled system can be divided into two basic phases: air and solidifying composite matrix (alloy) and dispersed phase of reinforcement. The air and composite are two immiscible phases between which appears a free surface. To describe this system we have implemented a well-known Volume of Fluid approach (VOF) which is based on the Euler frame. This approach also allows to consider surface tension forces in numeric models and enables to specify the interface between the phases on the basis of solving the continue equation for the volume fraction for one of the phase [10,11]. To model the


dispersed phase we have applied the Discrete Phase Model (DPM) which is based on the Lagrange frame. In addition, we have used the possibility of Fluent to a combination of these methodologies. Fluent allows to create Euler-Lagrange approach, which enables to take into account the interactions between the liquid matrix phase and the reinforcement particles. In the case of reinforcement particles we can define a set of forces that may affect the position of the particle. Fluent makes it possible to model solidification process by using the enthalpy method [9,12]. This method provides the manner of the modification of both transport and energy equations so that they can describe physical values of liquid and solid phases in the calculating domain.

The modeling of the casting process requires the selection of the appropriate calculation parameters to reflect the real system. The important element which may largely lead to differences between the results obtained from numerical simulations and experiment is the set of the parameters assumed in the calculations like the mass of the cast, physical and thermal properties of the material and thermal boundary conditions, etc. In the case of difference in simulation and experimental masses of the cast may considerably affect the time of composite solidification. Therefore, the main attention will be focused on the comparison of the course of the solidification process, the temperature distribution and final arrangement of the reinforcement particles but not on the solidification time.

Experimental casting

The experimental casting has been carried out in MMC Laboratory at the Silesian University of Technology. We have fabricated the composite suspension by the stirring method, described in detail in paper [1-5]. As the matrix material we have used a casting alloy of aluminum AlSi12CuNiMg (AK12), modified with a 2% magnesium and 0.03% strontium addition. Because of the objective of the paper we have applied two types of reinforcement particles (two types of composites): silicon carbide (SiC) of 10% mass fraction and a 50 µm particle size (AK12/SiC) and glass carbon Cg of 10% mass fraction and a 100 µm particle size (AK12/Cg). We have made all the castings by gravity casting into sand mold. For comparison, we have conducted an additional casting for the pure alloy AK12 (without reinforcement).

Fig. 1. Experimental system: (a) ThermaCAMTME25, (b) sand mold, point p1 indicates the localization of quick cup (thermocouple K).

The course of the solidification process we have recorded by means of the experimental system which enabled continuous control and measurement of the metal temperature during the solidification (Fig. 1). The solidification process of the matrix and AK12/SiC composite suspension has been recorded by using the ThermaCAMTME25 photometer system for temperature control and measurement. The point p1 (Fig. 1) at which the measurement has been performed, we have regarded in the calculation as the point p1. The system, equipped with a thermovision camera, LCD display and a laser pointer, has been connected to a SPIDER 8 recorder when used to monitor,


record and, simultaneously, visualize the temperature changes which take place during the solidification of the composite (Fig. 1). During the tests we have reordered the temperature and time of composite solidification as well as the mold temperature [1-5].

Simulation model

In order to carry out calculations which allow investigating the solidification process of alloy AK12, composites AK12/SiC and AK12/Cg, performed experimentally, we have used the two-dimensional model presented in Fig. 2. The simulation domain we have created by using unstructured rectangular mash. It consists of the following elements: the steel mount on which the sand mold is located and the fragment of steel pouring vessel. The outer border of the system we have defined by the pressure-inlet.

Fig. 2. 2D simulation model.

We have assumed that two continuous phases: air and liquid composite matrix exist in the designed model. We have used VOF approach to model the two-phase system mentioned earlier, air - composite matrix free surface, the volume fraction of particular continuous phases as well as the surface tension between the phases and the wall adhesion of the fluid on the solid surface [9-11]. This method introduces the parameter called volume fraction of k-phase fk which describes participation of k-phase in each calculation cell [9]:

( )

<<

=

interface fluid kat

fluid k inside

fluid k outside

1f0

1

0

t,f

th

th

th

k

k r (1)

According to the local value of fk, specific properties and variables are assigned to each control volume within the domain. In the calculation we have also taken into account the continuum surface force scheme (CSF) to model the surface tension in VOF calculations and the determination of the pressure jump across the interface [13].

To model the solidification of the composite matrix we have applied the enthalpy method [9,12]. This method assumes that the same system of differential equations and boundary conditions is counted in the entire computational domain and introduces the characteristic parameter called liquid fraction β. This parameter characterizes the state of material in specified point of the domain [9,12]:


>

<<−

−<

=β

L

LSSL

S

S

TTfor1

TTTforTT

TTTTfor0

(4)

where TL and TS are liquidus and solidus temperature, respectively. The value of liquid fraction parameter changes from 0 for solid phase to 1 for liquid phase. In the phase transition range (TS < T < TL) the liquid fraction assumes fractional value. The additional source term which is responsible for modification of energy and mass transfer equations during solidification process is described by the change of enthalpy of the material during this phase transition.

The presence of the dispersed phase (reinforcement particles) we have modeled using Discrete Phase Model (DPM). DPM assumes that the particles interact with continuous phase through a set of laws which are connected with the transfer of momentum, heat and mass [7-9]. In the carried out simulations for gravitational casting of the AK12/SiC and AK12/Cg the trajectories of individual particles can be treated as the balancing of the forces acting on them [7-9]:

bgpdp FF)uu(F

dt

du++−= (5)

where u and up are the fluid phase velocity and the particle velocity, respectively. Fd(u – up) is the drag force expressed as:

24

ReC

d

18F D

2pp

d ρ

µ= (6)

where µ is the viscosity of the fluid, CD is the drag coefficient [14], Re is the relative Reynolds number and dp is the particle diameter. Fg is gravitational force expressed as:

ppxg VgF ρ= (7)

where gx is gravitational acceleration, ρp is the density of the particle and Vd is volume of particle. Fb is buoyancy force expressed as:

pxb VgF ρ= (8)

where ρ is the density of the fluid. In our investigations we have assumed the model in which the reinforcement particles interact

with the continuous phase by a number of laws which describe the transfer of momentum, heat and mass. The momentum exchange appears as a momentum sink in the continuous phase momentum balance in any subsequent calculations of the continuous phase flow field whereas the heat transfer from the continuous phase to the discrete phase is computed by examining the change in thermal energy of a particle passing through each control volume in the model.

We have also assumed that viscosity of the composite matrix depends on the temperature according to linear relationship:

( )

>µ

<<µ+−−

µ−µ<µ

=µ

LL

LSSSSL

SL

SS

TTfor

TTTforTTTT

TTfor

(9)

where µS and µL are the viscosity of the solid and liquid of the composite matrix, respectively. The viscosity changes during the solidification, especially in the temperature range between TL and TS. Additionally, we have taken into account the presence of reinforcement particles and their effect on the matrix viscosity. The modification of viscosity has been performed by the equation [6]:

( )2ffpop V6.7V5.21 ++µ=µ (10)

where Vf is the volume fraction of particles.


Simulation parameters

We have assumed that the simulation domain possesses the outside boundary conditions as follows (pressure-inlet):

( )t,T)t,(T AAArr = (11)

For the mold wall and inside part of mount wall, the heat transfer is calculated directly from the solution in the adjacent cells:

( )lA TTq −α= (12)

where q is the heat flux, α is the fluid-side local heat transfer coefficient, TA and Tl are the wall surface and the local fluid temperatures, respectively. The fluid-side heat transfer coefficient is calculated on the grounds of the local flow-field conditions. We have assumed that the heat transfer on the interface between separated regions corresponds to the ideal contact condition as follows:

lATT ∇λ=∇λ nn (11)

where λ is the thermal conductivity of the solid, n is a normal to the wall.

Initial conditions and parameters

In order to carry out the simulations we have established the following parameters: pouring temperature of liquid composite 720ºC, initial temperature of the system 25ºC, outside border temperature (pressure-inlet) 25ºC, mass fraction of SiC and Cg 10%, diameter (spherical particles) 50 µm and 100 µm, respectively. Tables 1, 2 and 3 show the physical parameters of the all applied materials. We have assumed that the masses of the casting alloy AK12, composites AK12/SiC and AK12/Cg in the simulations refer to the casting masses 0.1341 Kg, 0.1339 kg (mass of the matrix 0.1205 kg, mass of SiC 0.0134 kg) and 0.1347 kg (mass of the matrix 0.1212 kg, mass of Cg 0.0135 kg) Due to the type of methods used in the calculation (methodology of the Eulerian-Lagrangian technique) the composite matrix mass has been the same in the calculation for the composite and alloy AK12.

We have assumed that the calculations have started from the initial state created during time-independent steady simulation. In order to obtain the initial temperature distribution the energy equation has been solved only for the system with two phases: air and liquid matrix in the selected location. The main time-dependent calculations have included additionally the counting of the flow and volume fraction equations.

Table 1. Physical properties of composite matrix AK12 and composites: AK12/SiC, AK12/Cg, . Values of parameters are based on Fluent database [9].

Material parameters Alloy AK12 AK12/SiC AK12/Cg density (liquid and solid) ρ heat capacity cp thermal conductivity λ viscosity µ latent heat L tension air-alloy interface γ contact angle

2680 kg/m3 981 J/kg·K 134 W/m·K 1.5×10-3 – 10.0 Pa·s 395 kJ/K 0.98 N/m2 120°

324 kJ/K

324 kJ/K

Table 2. Solidification parameters of composite matrix AK12 and composites: AK12/SiC, AK12/Cg, obtained from the experiment.

Material parameters Alloy AK12 AK12/SiC AK12/Cg solidus temperature TS liquidus temperature TL

559 °C (832 K) 572 °C (845 K)

555 °C (828 K) 555 °C (828 K)

551 °C (824 K) 553 °C (826 K)


Table 3. Physical properties of sand (mold), SiC and Cg (reinforcement). Values of parameters are based on Fluent database [9].

Material parameters Sand mold SiC Cg density (liquid and solid) ρ heat capacity cp thermal conductivity λ

1600 kg/m3 732 J/kg·K 0.825 W/m·K

3200 kg/m3 1010 J/kg·K 84 W/m·K

1800 kg/m3 1340 J/kg·K 150 W/m·K


Experimental course of temperature change. First, we compare the experimental data. The solidification process of matrix composite alloy AK12 as well as composites AK12/SiC and AK12/Cg has been presented in Fig. 3 by the solidification curves showing the change of the local temperature at the selected point p1 (Fig. 1) in the experiment.

Fig. 3 Change of local temperature vs. time for: AK12 (grey solid line), AK12/SiC (black solid line) and AK12/Cg (grey dashed line), AK12/SiC (black dashed line) vs. time obtained experimentally.

In Fig. 3 we can see the differences between the course of the solidification curves for the composite matrix (alloy AK12) and for each composite (AK12/SiC, AK12/Cg). These differences relate both to time and temperature at the start of the solidification. AK12 solidifies in the temperature range 572 – 559 °C during the 120 s. The temperature at the start of the solidification of the composite AK12/Cg is 553 °C, the composite solidifies in the temperature range 553 – 551 °C during 189 s. The fastest process of the heat exchange with the environment in the analyzed group of materials occurs for the composite AK12/SiC. The registered solidification time of the composite AK12/SiC is 52 s at the temperature 555 °C. Experimentally determined values of temperatures have been used in the simulations as the solidus and liquidus temperatures. The summary of these data are in Table 2.

Based on the presented data and Fig. 3 we can also indicate the differences of the cooling time

and the temperature range. The course of the solidification process depends on the presence or absence of reinforcement particles and their physical properties. The analysis of the result shows that solidification curves for AK12 and AK12/SiC have a similar course unlike AK12/Cg. A significant increase of the cooling time is the result of the physical properties of the glass carbon such as heat capacity.


Simulation course of temperature change. Fig. 4, 5 and 6 present the changes of the local temperature at the selected point p1 (black solid line) and maximum temperature in the system (black dashed line) during the simulation of the casting process for AK12, AK12/SiC and AK12/Cg, respectively. The change of the local temperature at the point p1 (Fig. 2) corresponds to the measuring point of the local temperature change during the experiment (Fig.1 – point p1). Additionally, for comparison, the figures include the appropriate experimental data (grey solid line).

Fig. 4 Change of local temperature at the point p1 (black solid line) and maximum temperature in the entire domain (black dashed line) vs. time for AK12 obtained during simulations. The grey solid line refers to change of temperature recorded experimentally. T1 and t1 are temperature and time at which solidification begins at the point p1, while T2 and t2, in which the solidification is completed.

Fig. 5 Change of temperatures vs. time for AK12/SiC – description like in Fig.4


Fig. 6 Change of temperatures vs. time for AK12/Cg – description like in Fig.4

The analysis of the Fig. 4, 5 and 6 indicates that the temperature change takes on the characteristic course (plateau) in the modeling system which arises from phase transition heat. If in any part of casting material, the temperature is in the range from Tliquidus to Tsolidus, the additional source of the energy appears and the phase transition takes place. Comparing the numerical and experimental results, we have noticed the discrepancies among the shapes of the temperature change curves at the point p1. The differences can arise from the assumed model of solidification which, as the literature shows, possesses several simplifications and physical parameters of the materials used in simulations. Besides, during the solidification process there appear some phenomena such as, overcooling of the casting, real changes in viscosity of individual components of the composite, which haven’t been included in the description of the solidification process. However, from analysis of figures we can see the similarity with regard to total time of the solidification of entire cast. The total solidification times for the assumed mass of the composite matrix AK12 and composites AK12/SiC and AK12/Cg in the simulation system are 335 s, 212 s and 241 s, respectively. Situation is different with regard to the start and end of the solidification time at the selected point p1. The solidification times determined in simulation and measured experimentally are presented in Table 4.

Table 4. Solidification time of composite matrix AK12 and composites: AK12/SiC, AK12/Cg, obtained from experiment and simulations at the selected point p1.

Time Alloy AK12 AK12/SiC AK12/Cg experiment simulation

70 s 241 s

52 s 146 s

189 s 170 s

Arrangement of reinforcement particles. The Fig. 7 and 8 show the final arrangement of the reinforcement particles of composites AK12/SiC and AK12/Cg, respectively, obtained during the simulation (a) and the experiment (b). The gray points represent the current position of individual parcels. For both cases, it is clear that the homogeneous distribution of the reinforcement particles in the pouring liquid cast substantially changes at the initial stage of the process. For the AK12/SiC, the SiC particles (larger density vs. matrix density) move down under the influence of the gravitational force (sedimentation) and form the distinct layer – Fig. 7a. The exactly opposite situation occurs for the system AK12/Cg. The Cg particles also form a distinct layer, but under the


influence of buoyancy force (lower density vs. matrix density), this layer is created in the upper part of the cast – Fig 8a. Thus, the effect of balancing the forces and the interaction of particles with the composite matrix as well as long-time cooling of the cast favor the accumulation of the particles. The simulation analysis of the particle arrangement we have based on DPM Concentration parameter. This parameter defines the total concentration of the discrete phase based on the unit quantity of density. The calculations show that for the assumed model, the maximum concentration does not exceed 2200 kg/m3 and 1200 kg/m3 for the systems AK12/SiC and AK12/Cg, respectively. We have observed a similar distribution of reinforcement particles in the cross-section for the castings of AK12/SiC (Fig. 7b) and AK12/Cg (Fig.8b).

Fig. 7. Final DPM Concentration for AK12/SiC obtained in simulational (a) and experimental (b) arrangement of reinforcement particles SiC. The gray points at (a) represent the position of individual parcels.

Fig. 8. Like in Fig. 7 but for AK12/Cg and reinforcement particles Cg.

Conclusions

One of the most important elements of the carried out simulations is the verification of the calculation possibility of the set of methods and models in ANSYS Fluent on the basis of experimental data. It can be concluded that the modeling the casting process of the metal matrix composite reinforced with ceramic particles is available in this program, especially arrangement of reinforcement. It is associated with selecting of the proper calculation parameters and taking into account the experimental data as well as creating of an appropriate simulation domain. The applied


theories and methods, however, contain a number of simplifications, which lead to the differences between simulation and experimental results. The next cause of the differences may result from the use of the two-dimensional system.

The course of the solidification process, shown by the solidification curve defined as the change of the local temperature vs. time, is closely related to the presence of reinforcement particles, their physical properties and the amount for both experiment and simulation. In both cases a long solidification time allows the formation of distinct layers of reinforcement particles. Therefore, the simulation enables the prediction of the final arrangement of reinforcement particles based on experimental data.

Acknowledgements

The present work is supported by the Ministry of Science and Higher Education grant PBU 77/RM4/2009 and National Science Center, Project No. N N508 630 540

References

[1] J. Śleziona, M. Dyzia, J. Wieczorek: Casting properties of composites suspensions AlSi-SiC, Archiwum Odlewnictwa Rok 2006, Rocznik 6, Nr 22, Archives of Foundry Year 2006, Vol. 6, N0 22, PAN- Katowice PL ISSN 1642-5308 (in polish).

[2] A. Dolata-Grosz, M. Dyzia, J. Śleziona: Solidification and structure of heterophase composite, Journal of Achivements in Materials and Manufacturing Engineering, 20, (2007) 103-106.

[3] A. Dolata-Grosz, M. Dyzia, J. Śleziona, Solidification curves and structure of heterophase composites, Archives of Materials Science and Engineering, 31, (2008) 10-15.

[4] A. Dolata-Grosz, M. Dyzia, J. Śleziona, The formation of the structure of cast composites in different solidification conditions, Archives of Materials Science and Engineering, 31 (2008) 13-16.

[5] A. Dolata-Grosz, M. Dyzia, J. Śleziona, Solidification analysis of AMMCs with ceramic particles, Archives of Materials Science and Engineering, 28 (2007) 401-404.

[6] J. Sobczak, Metal Matrix Composites (Kompozyty Metalowe), Instytut Odlewnictwa i Instytut Transportu Samochodowego, Kraków – Warszawa 2001 (in polish).

[7] R. Zagórski, Implementation of computer simulation for modeling arrangement of ceramic reinforcing particles during casting process of metal matrix composite, Int. J Mater From, 3 (2010) 655-658.

[8] R. Zagórski, J. Śleziona, Influence of thermal boundary condition on casting process of metal matrix composite, Archives of Materials Science and Engineering, 42 (2010) 53-61.

[9] www.ansys.com

[10] C.W. Hirt, B.D. Nichols, Volume of fluid (VOF) method for the dynamics of free boundaries, J. Comput. Physics, 39 (1981) 201-225.

[11] J.U. Brackbill, D.B Kothe, C. Zemach, A continuum method for modeling surface tension, J. Comput. Physics, 100 (1992) 335-354.

[12] V.R. Voller, M. Cross, N.C. Markatos, An enthalpy method for convection-diffusion phase change, Int. J. Num. Meth. Eng. 24 (1987) 271-284.

[13] J. U. Brackbill, D.B. Kothe, C. Zemach. A Continuum Method for Modeling Surface Tension, J. Comput. Phys.. 100 (1992) 335–354.

[14] A. Morsi, A.J. Alexander , An Investigation of particle trajectories in two-phase flow systems. J. Fluid Mech., 55 (1972) 193-208.


CHAPTER 2:

Magnesium and Magnesium Alloys

Plasticity and microstructure of hot deformed magnesium alloy AZ61

Dariusz Kuc 1,a, Eugeniusz Hadasik 1,b,Iwona Bednarczyk 1,c.



Keywords: AZ61 magnesium alloy, plastometric tests, hot compression, Zener-Hollomon parameter, flow stress, microstructure

Abstract. The article presents the results of tests connected with the influence of strain parameters on the change of flow stress and microstructure of magnesium alloy AZ61 (symbol according to ASTM norms). Test of uniaxial hot compression were conducted in temperature range from 250 to 400°C and the strain strain rate from 0.01 to 1 s-1. Analysis of plastometric tests and microstructure observation allowed to establish which mechanism - slip or twinning – is dominant in particular conditions of shaping AZ61 alloy. Achieved results were compared with previous results achieved for AZ31 alloy type with lower content of aluminium.

Introduction

The current trends in the automotive and aviation industry focus mainly on reduction of the vehicle weight and on saving energy (particularly on saving fuel) and thereby protecting the environment [1,2]. Such a set of technical, economical and ecological aspects arouses a considerable interest of the industry in light alloys. Owing to a number of their advantageous mechanical properties including, first of all, low density (1.74 g/cm3), magnesium alloys are more and more frequently used as an engineering material. Additional beneficial properties of magnesium alloys, including the ability to suppress vibrations, cause more and more often application of those alloys as elements for construction of electronic items, in ballistics and in production of sports equipment [2-4]. Casting processes are still most often applied for the production of components from magnesium alloys, due to very good casting properties of magnesium. Alloys used for plastic working are less popular compared to those processed via casting and therefore, the number of their grades is much smaller. The number of alloying components in cast magnesium alloys is always higher than in alloys subject to plastic working. Alloys from the group Mg-Al-Zn-Mn have the best set of properties, as they contain up to 8% Al with an addition of Mn (up to 2%) and Zn (up to 1.5%) [2-5]. Among the elements subjected to plastic working, metal sheets deserve special attention as they can be applied for the construction of light vehicles [2-6]. Due to the beneficial properties of plastically deformed magnesium alloys there is a tendency to increase its production. It is very important for the scientific and technological background kept testing the new and better ways of application of light alloys for construction elements [4]. Conducted works, due to the complexity of the phenomena occurring in microstructure of Mg-Al-Zn alloys, should concentrate on showing the strain mechanisms and processes of rebuilding structure during strain [6-10] and the influence of heat treatment on its properties [11]. In the process of plastic deformation it is also necessary to analyse the course of the structural phenomena occurring in the break between strains. The basic aspect deciding on the presence of the influence of the preceding effects on the next stages of shaping is the duration of the break between such strains. It allows for representation of the real conditions of the shaping process and helps in improvements of plastic treatment technology to achieve improved utility properties [6, 9-11]. The paper presents the tests which aim at the assessment of the structure changes during heating and hot deformation of magnesium alloy with chemical composition of the type AZ61 according to ASTM norms. Magnesium alloy AZ61 is processed with the use of extrusion and hot forging. At present, works are conducted to check the possibilities of rolling the alloy. The advantages of AZ61


alloy are the more beneficial resistance properties in comparison to AZ31 type. The main disadvantages of AZ61 alloy are the occurring fusible eutectics in temperature of about 450°C, which limit temperature range of hot plastic treatment. Plasticity tests were conducted with the use of uni-axial hot compression in temperature range 250÷400°C and strain rate 0.01÷1s-1. The paper presents relationships between characteristics of plasticity such as: maximum flow stress (σpp) and strain corresponding to maximum and the Zener-Hollomon parameter. Characteristics of plasticity and microstructure of AZ61 alloy were compared to the behaviour of AZ31 alloy, with lower aluminium content which was tested earlier [9,10]. Achieved plasticity characteristics and changes of structure during continuous deformation will be applied in elaboration of complex mathematical model of structure changes of AZ61 alloy during high-temperature strain.

Material for tests

Materials for tests were extrusion rods with diameter of 12mm from AZ31 alloy (according to ASTM) and chemical composition (%mass): 6.10% Al, 0.95% Zn, 0.6% Mn, 0.002% Cu. Rods after extrusion were annealed at temperature of 400oC with soaking time of 40 minutes and next cooling on air. Microstructure of the alloy in condition after extrusion and annealing is shown in Fig. 1. In the initial state after extrusion the structure of the alloy is fine-grained (Fig. 1a), and after annealing the alloys shows a tendency to grain growth (Fig. 1b).

a)

b)

Fig. 1. Microstructure AZ61 alloy: a) after extrusion forging, b) after annealing at 400oC/40min, with air cooling

X- ray diffractograms for AZ61 magnesium alloy and the previously tested AZ31 alloy are presented in figure 2. After conducted heat treatment the presence of solid solution of α-Mg was found. On the diffractogram of AZ61 alloy the diffraction lines shift becomes visible (Fig. 2b). It proves the change of the elementary magnesium cell size resulting from the increase of the aluminium in the solid solution of α-Mg.


a)

co

un

ts

2θθθθ b)

co

un

ts

2θθθθ

Fig. 2. X-ray pattern of investigated magnesium alloys after annealing at temperature of 400ºC: a) AZ31, b) AZ61

Methodology of tests

Tests of hot compression were conducted using heat mechanical simulator Gleeble 3800. Samples were heated to a temperature of 400°C with strain rate of 3°C/s, holding in that temperature for 300 s, and then cooled with the strain rate of 5°C/s to test temperature (250, 300, 350 or 400°C). Strain was initiated after 30s. The following strain rate were applied: 0,01, 0,1, 1,0 s-1 with the real rolling reduction equal 1.0. During compression tests from the strain strain rate of 1 s-1 the increase and then the later drop of temperature of compressed samples was observed. For smaller strain rate some slight deviations of the set temperature of samples were observed only in the initial phase of compression. The result of temperature increase is bigger when the initial temperature of the test is lower. The measurements show that the system of temperature regulation in the simulator is able to keep a constant temperature of the samples only during compression at a lower strain rate. Taking into consideration the temperature changes of the compressed samples the dependencies of stress from strain achieved in each of the tests were corrected. In order to determine the value of flow stress in complex temperatures the values of stress achieved in various temperatures for the same strain values and strain strain rate as well as the method of linear approximation were used. Structural assessment was conducted with the use of light microscope by “Olympus” company, in bright field technique.


Results of investigations

The influence of temperature and strain rate of strain on the chosen flow curves of AZ61 alloy are shown in figures 3 and 4. Similarly to AZ31 alloy tested earlier [10], the decrease of compression temperature led to the change of the flow curve. It signifies the increase of the mechanical twinning influence as the the main mechanism of plastic deformation (Fig. 3). Similar behaviour of the alloy is observed with the increase of strain rate (Fig. 4). By bigger of strain rate and lower temperatures the flow curves show a significant growth of flow stress for small values of strain and then a sudden drop to fixed value. A characteristic aspect of those curves is the fact that the maximum of flow stress is achieved for similar value of strain, about 0.2. By lower strain rate of strain and higher temperatures the change of values of flow stress is much smoother and the strain necessary to achieve the maximum value of flow stress increases in linear way together with the decrease of temperature of the shaped material. Distinctness of the alloy behaviour depending on the parameters of the process is confirmed by microstructure tests after compression for the strain value of ε = 0.1, for temperatures of 250°C and 400°C (Fig. 5a, b). In microstructure after compression in temperature of 250°C the process of reconstruction of the structure as a result of dynamic recrystallization is preceded by appearance of a big amount of twins (Fig. 5a). In higher temperatures, an intensive migration of crystalline boundaries is observed by slight increase of frequency of twinning (Fig. 5b). Changes of the maximum flow stress (σpp) and the corresponding strain εp in function logZ is shown in Fig. 6 and 7. Zener-Hollomon parameter (Z) was calculated with the use of dependency:

expQ

ZR T

ε = ⋅ (1)

on the basis of activation energy Q marked in the program ENERGY 3.0 [12] according to a constitutive equation:

exp (sinh( )n

pp

QC

R Tε ασ

= − × ⋅ (2)

where: C [s-1], α [MPa-1], n [-] – coefficients Energy of alloy activation is small and equals 171.3 kJ/mol. The value of energy is slightly lower from the value for alloy AZ31 – 175.1 kJ/mol (which was marked earlier). This value means that there is a little influence of temperature on the flow stress and the tendency to dynamic recrystallization process of the material with small arrangement defect energy of which is decreased with the increase of the aluminium amount respectively for alloy AZ31 and AZ61to 27.8 mJ/m2 and 16.4 mJ/m2 [7]. A dependency of power character between the maximum flow stress (σpp) and Zener-Hollomon parameter (Z) was found. In the whole tested range of parameter variation the alloy AZ61 has higher values of maximum flow stress (Fig. 6) With the same parameters of the process the strain value εp is higher for alloy AZ31. For classic course of flow curve a dependency of power character in Z parameter function was found, whereas for parameter range where twinning dominates such dependency is stable and the value of εp is close to 0.2 (Fig. 7). The course of the flow curves suggests that is conditions of strain for which the Zener-Hollomon parameter is lower than 4,13×1015 s-1, the dominant mechanism of plastic strain is slip. For bigger values of Z parameter twinning begins to dominate. Similar dependency was found for AZ31 alloy.


0

50

100

150

200

250

0 0,1 0,2 0,3 0,4 0,5 0,6 0,7 0,8 0,9 1 1,1 1,2

Strain εεεε

Flo

w s

tres

s σσ σσ

p, M

Pa

350°C

300°C

250°C

400°C

Fig. 3. Flow stress curves for AZ61 alloy after deformation at temperature range of 250°C to 400°C

with a rate of 0.1 s-1

0

50

100

150

200

250

300

350

0 0,1 0,2 0,3 0,4 0,5 0,6 0,7 0,8 0,9 1 1,1 1,2

Strain εεεε

Flo

w s

tre

ss σσ σσ

p, M

Pa

0,01s-1

0,1s-1

1s-1

Fig. 4. Flow stress curves for AZ61 alloy after deformation at 250°C with a rates of 0.01, 0.1 s-1

and 10 s-1

a) b)

Fig. 5. Microstructure of AZ61 alloy after deformation at temperature: a) 250°C, b) 400°C with a rate of 1s-1

migration GB

twinning


0

50

100

150

200

250

300

350

24 26 28 30 32 34 36 38 40 42

Zenera Hollomon parameter, Log Z

Pe

ak f

low

str

es

s σσ σσ

pp

, M

Pa

AZ31

AZ61

σσσσpp= 11,7×105

Z0,127

σσσσpp= 11,6×105

Z0,141

Fig. 6. Peak flow stress (σpp) as a function of the log Z, Z - Zener-Hollomon parameter

0

0,05

0,1

0,15

0,2

0,25

24 26 28 30 32 34 36 38 40 42

Zener Hollomon parameter, Log Z

Str

ain

εε εεp

AZ31

AZ61

Fig. 7. Deformation εp corresponding to the maximum flow stress σpp as a function of the Zener-Hollomon parameter

The influence of temperature and strain rate of compression to strain ε = 1 on the microstructure of the tested alloy is shown is Fig. 8. After compression in temperature of 250°C for both applied strain strain rate 0.01 and 1s-1 the structure of primary elongated crystallites and ultra-fine dynamically recrystallized grains was observed (Fig. 8a, 8b). Samples which are deformed with smaller strain rate are characterised with bigger advancement of recrystallization process, recrystallized grains are observed both on the boundaries and in the area of primary crystallites (Fig. 8a). After strain at temperature of 300°C and 350°C with strain rate of 0.01s-1 the microstructure consists of fine dynamically recrystallized grains (Fig. 8c, 8e), and for the bigger strain rate some non-recrystallized areas are observed surrounded by small chain of new crystallites (Fig. 8d and 8f). The growth of crystallite size becomes visible with increase of temperature of compression process. At temperature of 400°C the recrystallization process intensifies and microstructure is fully recrystallized for both applied strain rates of compression (Fig. 8g, 8h).


Fig. 8. Microstructure of AZ61 alloy after deformation at temperature range from 250°C to 400°C

with a rate: of 0,01s-1, b) 1s-1. Strain ε = 0.1. Summary The paper deals with description of tests of influence of strain parameters on plasticity and structure of magnesium alloy AZ61 (Mg-Al-Zn). Conducted tests with the use of Gleeble 3800 simulator with the use of uni-axial compression method enabled to mark the flow curves in the system stress -strain of the tested alloy. In order to mark the values of flow stress (σpp), more precisely, a correction of the values due to temperatures of the process was conducted. For the assessment of the sample structure the samples were intensively cooled with water for the so-called freezing of the structure. Alloy AZ61 is characterised with mono-phase structure which is confirmed by conducted X-ray tests. Achieved X-ray diffactograms do not show the presence of non-dissolved phases (Fig. 2). For the applied temperature range and strain rate of the process, which covers the range of the parameters of plastic treatment, no cracks were found in compressed samples. Flow curves depending on the parameters of strain show two different mechanisms of strain (Fig. 3, 4). For higher temperatures and lower strain rate of strain the curve has classic course of flow stress changes. In lower temperatures and higher strain rate of strain the course of stress changes is different and characteristic for the twinning process which is confirmed by structural tests for small strain values (Fig. 5). Similar dependencies were achieved for alloy AZ31. It was proved here that there is a strong dependency of power character between maximum flow stress (σpp) and Zener


– Hollomon parameter (Z). Alloy AZ61 in the whole tested range of strain parameters variation shows bigger values of flow stress than alloy AZ31 (Fig. 6). It results from bigger consolidation of the solid solution by two times bigger content of aluminium in tested material. Strain corresponding to maximum flow stress εp changes exponentially with increase of Z parameter. For conditions of the process corresponding to twinning, value of strain εp is stable (about 0.2) and independent from the changes in compression parameters. Conducted tests of microstructure show the process of dynamic re-crystallisation in the whole tested range of process parameters. The presence and the size of re-crystallised grains is dependent on temperature and strain rate. In temperature of 250°C, the bimodal microstructure is achieved, constructed of elongated primary crystallites and ultra –fine recrystallized grains which appear mainly in the area of primary grain boundaries (Fig. 8a, b). Fully recrystallized microstructure was observed after compression in temperature of 300°C with strain rate of 0.01s-1 (Fig. 8c). Increase of strain temperature leads to size growth of recrystallized grain (Fig. 8e, g). Achieved results will be used to elaborate a complex model of structure changes of hot deformed magnesium alloy AZ61. However, there is a need to take into account in the model the varied mechanisms of deformation in analysed alloy depending on the parameters of the process which will allow for correct planning of rolling and extrusion technology of the products made out of this alloy.

Paper conducted within Project “Modern material technologies applied in aviation industry”, No. POIG.0101.02-00-015/08 in Operational Program Innovative Economy (POIG). Project co-financed by European Union from the funds of European Fund of Regional Development.

References

[1] H. Friedrich, S. Schumann, Research for a „new age of magnesium“ in the automotive industry, Journal of Materials Processing Technology. 117 (2001) 276÷281.

[2] R. Kawalla, Magnesium and magnesium alloys Monograph edited by Hadasik E., Manufactured of metals. Plasticity and structure. Silesian University and Technology, Gliwice, 2006.

[3] B. L. Mordike, T. Ebert, Magnesium Properties - applications – potential, Materials Science and Engineering. A302 (2001) 37÷45.

[4] J. Bohlen, D. Letzig K. Kainer U, New Perspectives for Wrought Magnesium Alloys, Materials Science Forum. 546-549 (2007) 1÷10.

[5] E. Hadasik., D. Kuc, G. Niewielski, R. Śliwa, Development of magnesium alloys for plastic working, Hutnik - Wiadomości Hutnicze 76. 8 (2009) 666÷670.

[6] B. Jiang , J. Wang, P Ding, Ch Yang, Rolling of AZ31 Magnesium Alloy Thin Strip Materials Science Forum. 546-549 (2007) 365÷368.

[7] H. Somekawa, Dislocation creep behaviour in Mg-Al-Zn alloys, Materials Science and Engineering A. 407(2005) 53÷61.

[8] M. M. Myshlyaev, H. J. McQueen, E. Konopleva, Microstructural development in Mg alloy AZ31 during hot working, Materials Science and Engineering A337 (2002) 121÷127.

[9] D. Kuc , E. Hadasik, G. , A. Płachta, Structure and plasticity of the AZ31 magnesium alloy after hot deformation, Journal of Achievements in Materials and Manufacturing Engineering 27(2008) 27÷31.

[10] D. Kuc, E. Hadasik, A. Szuła, Research of plasticity and microstructure of magnesium alloys AZ31 type in die – casting and hot rolling condition after deformation, Hutnik - Wiadomości Hutnicze, 76. 8 (2009) 666÷670.

[11] L.A. Dobrzański, T. Tanski, L. Cizek, J. Madejski, The influence of the heat treatment on the microstructure and properties of Mg-Al-Zn based alloys, Archives of Materials Science and Engineering. 36/1 (2009) 48÷54.

[12] I. Schindler J. Boruta, Determinig and using of activation energy for hot working processes, Archives of Metalurgy 4. (1994) 471÷491.


Effect of modification on the structure and properties of QE22 and RZ5 magnesium alloys

Stanisław Roskosz1, a, Bartłomiej Dybowski1,b , Janusz Paśko2,c 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

2 Zakład Metalurgiczny „WSK Rzeszów”, ul. Hetmańska 120, 35-078 Rzeszów, Poland

[email protected], [email protected], [email protected]

Keywords: magnesium casting alloys, QE22 alloy, RZ5 alloy, microstructure, fractography, mechanical properties, quantitative metallography

Abstract. Magnesium alloys are the lightest, widely used structural material. They are often used in aeronautical and automotive industries, where the weight savings are essential. Magnesium alloys present acceptable mechanical properties but their high temperature properties are unsatisfactory. This led to development of magnesium alloys with rare earth elements addition. To achieve good mechanical properties these alloys are modified with zirconium. Modification affects positively also corrosion resistance of Mg-RE alloys. It is important to study impact of modifier amount on the structure and properties of these alloys. Unmodified and modified alloys were investigated. Three variants of modification were: modification according to Magnesium-Elektron (MEL) specification, 50% and 100% more modifier. Mechanical and structural properties were investigated. Fractures were observed on scanning electron microscope. Results showed that grain refinement and yield strength increase with increasing amount of modifier. Impact of modification on tensile strength is unclear, probably because of non-metallic inclusions in the material’s structure. The inclusions sources are oxygenated nappe of liquid metal and fluxes, used during smelting. Introduction

QE22 casting magnesium alloy with rare earth elements guarantee high mechanical properties up to 200˚C [1, 2]. It is caused by silver addition, which strongly increase response to age hardening [3, 4]. The main drawbacks of this alloy are: weak corrosion resistance and high cost [5]. RZ5 magnesium alloy with rare earth elements addition is characterized by high content of Zn. This alloy guarantee good casting properties and weldability, connected with acceptable mechanical properties [2]. Even though zinc in magnesium alloys with rare earth elements decrease their response to age hardening, it can be considered as a solid solution strengthening element [6]. Mg-RE alloys obtain demanded mechanical properties after grain size modification with zirconium [7]. The modification causes significant grain refinement and improves casting properties of the alloy [8]. Modification positively affects corrosion resistance of the Mg-RE alloys by solid solution stabilization and by creating more uniform grid of intermetallic phases, characterized by better corrosion resistance [9]. The article presents results of the studies on impact of amount of the modifier on grain refinement efficiency and alloy’s properties. Materials for the research

Two magnesium casting alloys with the addition of rare earths elements constituted the materials for testing: QE22 and RZ5, with no modification of the chemical composition and with 3 variants of modifications: according to the MEL specifications, +50% and +100%. The chemical compositions of tested alloys are presented in table 1. The samples for testing, made in Zakład Metalurgiczny "WSK Rzeszów", were prepared by gravity sand-casting. For every variant of modification, tensile strength test samples were cast, according to the PN-91/H-88052 norm and samples for linear shrinkage tests according to the DIN and TGL(103-2011) norms.


Table 1. Chemical composition of tested alloys (wt. %) Alloy Mg Zn Zr RE Ag Cu QE22 bal. - 0.6 2.0 2.5 0.07 RZ5 bal. 3.5-5.0 0.4-1.0 0.8-1.7 - -

Methodology

For every variant of modification, a static tension test on testing machine Zwick 1474 as well as impact strength test on Charpy impact test machine in room temperature were performed. On the fractures after the impact strength tests fractographic tests were conducted, using Hitachi S3400N scanning microscope (SEM). The examinations of the microstructure were carried out on the Olympus GX71 light microscope. The specimens were grinded on the abrasive papers with gradation 320, 500 and 1200 and polished on diamond pastes with grain equal 3µm and 1µm. Finishing polishing was performed on Al2O3 paste with grain size 0,25µm. Measurement of the eutectic areas were conducted on unetched sections, with 500x magnification. The grain size measurement was performed on etched sections with the etchant of the following chemical constitution: 4.2 g picric acid, 10 ml H2O, 10 ml CH3COOH, 70 ml C2H5OH. The detection and measurement of the components of the structure was conducted using Met-Ilo v. 12.1.

Results of the investigations

Mechanical examinations. The mechanical properties of the alloys are presented in table 2.

Table 2. Mechanical properties of the cast alloys, unmodified and after modification. Alloy Rm

[MPa] R0,2

[MPa] R0,2 / Rm

[-] A5

[%] U

[J/cm2] QE22 - unmod. 134 85 0.63 3.3 59.4 QE22 - mod. acc. to MEL 118 100 0.85 1.1 138.1 QE22 - mod. +50% 139 112 0.81 2.1 207.1 QE22 - mod. +100% 154 112 0.73 3.2 187.7 RZ5 - unmod. 134 88 0.66 2.2 85.8 RZ5 - mod. acc. to MEL 180 121 0.67 3.7 - RZ5 - mod. +50% 168 128 0.76 2.3 102.2 RZ5 - mod. +100% 156 127 0.81 1.4 121.5

Fractographic tests. The fractures both before and after the modification showed fragile, inter-crystallic character. On the surface of fractures, numerous non-metallic inclusions were found, containing Mg, alloy elements and Oxygen (fig. 1a). The fractures of the modified QE22 alloy alone showed no inclusions. (fig. 1b). On the surface of all the fractures, secondary crackings are present (fig. 1c), whereas on the fractures of QE22 alloy, numerous voids were found (fig. 1d). Structural examinations. The structure of the examined alloys consisted of a solid solution of alloy elements in magnesium α, as well as an eutectic created by inter-metallic phases Mg-RE-Ag (QE22) or Mg-RE (RZ5), and Mg α solid solution (fig. 2). The eutectics were emitted in inter-dendritic spaces.


a)

b)

c)

d)

Fig. 1. The surfaces of the tested fractures, SEM, SE, a) the surface of the fracture of the sample of the unmodified QE22 alloy, b) the surface of the fracture of the +50% modified QE22 alloy, c) secondary crackings on the surface of the fracture of the +50% modified RZ5 sample, d) voids on the surface of the QE22 +100% modified alloy

a)

b)

c)

d)

Fig. 2. The structure of tested alloys, LM. a) the eutectis emitted in the inter-dendritic spaces, QE22 alloy, unmodified, b) porosity in the QE22 sample modified according to MEL specifications, c) cellular structure of the grain, RZ5 alloy, modified according to MEL, d) the cellulo-dendritic structure of the grain, QE22 alloy, modified according to MEL.


The results of the quantitative analysis of the eutectic precipitates are presented in tab. 3, whereas the results of the analysis of the size and the shape of the grain in tab. 4 and 5.

Table 3. The results of the quantitative analysis of the eutectic precipitates Alloy Surface

fraction AA [%]

Variation coefficient υ(AA)[%]

Linear fraction LL [%]

Relative area of grain boundaries SV [µm2/ µm3]

QE22 - unmod. 5.18 21.4 5.31 0.56 QE22 - mod. acc. to MEL 6.90 15.1 6.73 0.587 QE22 - mod. +50% 6.10 14.3 5.99 0.664 QE22 - mod. +100% 5.85 15.2 5.92 0.662 RZ5 - unmod. 3.57 20.6 3.51 0.712 RZ5 - mod. acc. to MEL 4.48 8.9 4.45 0.766 RZ5 - mod. +50% 4.57 7.9 4.47 0.778 RZ5 - mod. +100% 4.32 10.7 4.36 0.772

The test results analysis

On the surface of the fractures, one could observe numerous non-metallic inclusions. These inclusions were of a twofold character: the first type impurities contained only alloy elements and the oxygen. These impurities were most likely included from the casting of the oxygenated nappe of liquid metal. The second type of impurities consisted of inclusions coming from fluxes used during smelting. Along with the increase of the amount of the modifier, monotonous decrease of the grain size occurs and the increase of the volume fraction of the eutectics (fig. 3). With the greater volume fraction of the eutectics, they tend to break up - the relative surface of the eutectics boundaries SV

increases (tab. 3).

Table 4. The results of the size measurements of the QE22 grain

Parameter symbol unit QE22 - unmod.

QE22 - mod. acc. to MEL

QE22 - mod. +50%

QE22 - mod.

+100% grain size

area of flat section A [µm2] 9229 2484 774 543 number of grain per unit area NA [mm-2] 106 395 1274 1808 relative area of grain boundary

SV [µm2/µm3] 0.026 0.048 0.084 0.105

heterogeneity of the grain size variation coefficient A ν(A) [%] 102 83 81 77

grain shape shape factor ξ - 0.635 0.662 0.662 0. 644 elongation factor δ - 1.64 1.67 1.78 1.65


Table 5. The results of the size measurements of the RZ5 grain Parameter symbol unit RZ5 -

unmod. RZ5 -

mod. acc. to MEL

RZ5 - mod. +50%

RZ5 - mod.

+100% grain size


SV [µm2/µm3] 0.029 0.084 0.103 0.110


grain shape shape factor ξ - 0.627 0. 582 0. 617 0. 605 elongation factor δ - 1.59 1.69 1.63 1.62 a)

b)

Fig. 3. a) the results of the mid area plane section measurements of the grain, as well as the volume fraction of the eutectics with the variant of modification for QE22 alloy, b) the results of the average area of plane section measurements of the grain, as well as the volume fraction of the eutectics with the variant of modification for RZ5 alloy.

In all the cases, with the increase of the amount of the modifier, the yield strength increased as well (fig. 4) and the proportion of the yield strength to the tensile strength (tab. 2). a)

b)

Fig. 4. The relation between tensile strength and yield strength and the modification variant:

a) QE22 alloy, b) RZ5 alloy.

0

2000

4000

6000

8000

10000

012345678

QE22 -

unmod.

QE22 -

mod.

acc. to

MEL

QE22 -

mod.

+50%

QE22 -

mod.

+100%

[µm

2]

[%]

eutectics volume fraction VV[%]

surface area of the grain plain section A [µm2]

010002000300040005000600070008000

0

1

2

3

4

5

RZ5 -

unmod.

RZ5 -

mod.

acc. to

MEL

RZ5 -

mod.

+50%

RZ5 -

mod.

+100%

[µm

2]

[%]

eutectics volume fraction VV[%]

surface area of the grain plain section A [µm2]

0

20

40

60

80

100

120

0

20

40

60

80

100

120

140

160

180

QE22 -

unmod.

QE22 -

mod. acc.

to MEL

QE22 -

mod.

+50%

QE22 -

mod.

+100%

[MP

a]

[MP

a]

Rm [Mpa] R0,2 [Mpa]

0

20

40

60

80

100

120

140

020406080

100120140160180200

RZ5 -

unmod.

RZ5 -

mod.

acc. to

MEL

RZ5 -

mod.

+50%

RZ5 -

mod.

+100%

[MP

a]

[MP

a]

Rm [Mpa] R0,2 [Mpa]


Further increase of the amount of modifiers causes insignificant increase of those properties. The influence of the amount of the modifier on tensile strength, elongation and impact strength is ambiguous. This results from the contamination of the tested alloys with non-metallic inclusions (fig. 2a). Conclusions

1. The modification of the QE22 and RZ5 alloys strongly influences the grain refinement, also

causing the increase of the eutectic areas and their refinement. 2. With the increase of the amount of the modifier the yield strength of the alloys increases as

well. It stays in concordance with the Hall-Petch relationship - the yield strength of the alloy is proportionate to the refinement of the grain.

3. The influence of the amount of the modifier on tensile strength, elongation and impact strength is ambiguous. This is caused by the presence of non-metallic inclusions, whose sources are: the oxygenated nappe of liquid metal and fluxes. Both types of impurities present errors in casting technology.

4. The modification of QE22 and RZ5 alloys according to MEL specification causes significant grain refinement. Further increase of the amount of the modifier does not increase its refinement and yield strength significantly. The set of the modified alloy properties according to MEL specification is good. From the economic perspective, there is no justification for further increase of the amount of modifiers.

Acknowledgment

The present work was supported by the Polish Ministry of Science and Higher Education under the research project No 6ZR7 2009C/07354.

References:

[1] I. Stloukal, J. Čermák, Silver diffusion in commercial QE22 magnesium alloy with Saffil fiber reinforcement, Composites Science and Technology 68 (2008) 2799–2803

[2] Magnesium Elektron, Magnesium casting alloys, Datasheet: 440 [3] T. Rzychoń, A. Kiełbus, J. Cwajna, J. Mizera, Microstructural stability and creep properties of

die casting Mg-4Al-4RE magnesium alloy, Materials Characterization 60 (2009) 1107-1113 [4] Y.M. Zhu, A.J. Morton, J.F. Nie, Improvement in the age-hardening response of Mg–Y–Zn

alloys by Ag additions, Scripta Materialia 58 (2008) 525–528 [5] A. Kiełbus, Microstructure and properties of sand casting magnesium alloys for elevated

temperature applications, Solid State Phenomena, 176 (2011) 63-74 [6] B. Bronfin, A. Ben-Dov, J. Townsend, S. Mahmood, J. Vainola, S. Deveneyi, N. Moscovitch,

Advanced gravity casting magnesium alloys for the aircraft industry, Magnesium, Edited by K.U. Kainer, WILEY-VCH (2007) 14-19

[7] G. Song, D. StJohn, The effect of zirconium grain refinement on the corrosion behaviour of magnesium-rare earth alloy MEZ, Journal of Light Metals 2 (2002) 1–16

[8] P. Lyon, I Syed, S. Haeney, Elektron 21 – An aerospace magnesium alloy for sand cast & investment cast applications, Magnesium, Edited by K.U. Kainer, WILEY-VCH (2007) 20-25

[9] M. Qian, D.H. StJohn, M.T. Frost, Heterogeneous nuclei size in magnesium–zirconium alloys, Scripta Materialia 50 (2004) 1115–1119


Influence of mould cooling rate on the microstructure of AZ91 magnesium alloy castings

Stanisław Roskosz 1, a, Bartłomiej Dybowski 1,b , Robert Jarosz 2,c 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

3 Zakład Metalurgiczny „WSK Rzeszów”, ul. Hetmańska 120, 35-078 Rzeszów, Poland


Keywords: magnesium casting alloys, AZ91 alloy, microstructure, quantitative metallography

Abstract. Magnesium alloys are the lightest, widely used structural material. They are often used in

aeronautical and automotive industries, where the weight savings are essential. Due to high

responsibility of the elements made from magnesium alloys it is important to achieve high quality

castings without any defects. The paper presents results of investigations on influence of sand

mould cooling rate on microstructure and quality of the castings. Six identical castings, fed and

cooled in different ways were investigated. Studies consisted of: RTG investigations and SEM and

LM observations. Microstructure was evaluated qualitatively and quantitatively. RTG investigations

showed that casting without feeder and cooler, casting only with feeder and castings cooled with

20mm and 40mm thick cooler contains voids inside. Castings with feeder and coolers 20mm and

400mm thick were flawless. Microstructure evaluation showed that castings with and without

defects have different structure. Castings with defects were characterized by higher volume fraction

of Mg17Al12 intermetallic phase. Flawless castings were characterized by fully divorced eutectic.

Introduction

AZ91 alloy is one of the most commonly used magnesium casting alloys, widely applied in aircraft,

automotive and electric industries, where the lowering of elements weight is important [1, 2]. The

equilibrium microstructure of the AZ91 alloy consist in 100% of Mg α solid solution [3]. Due to

strong segregation of alloying elements in liquid state in the first stage of solidification, solidifying

Mg solid solution contains about 3-4% of Al [4]. In the last stages of solidification (in eutectic

temperature), the remaining liquid alloy, rich in alloying elements rapidly solidify [4,5]. Eutectics in

AZ91 alloy have two forms: fully divorced – massive Mg17Al12 particles are surrounded by eutectic

α Mg or partially divorced – with eutectic α Mg “islands” within the Mg17Al12 particles and

partially surrounding them [6,7]. Moreover, during slow cooling rate, fine Mg17Al12 particles

precipitate from supersaturated solid solution [3].

Microporosity created during solidification affects strongly mechanical properties of AZ91 alloy.

With increasing volume fraction of pores linearly decrease yield strength. Elongation and tensile

strength decrease parabolic with increase of microporosity [8]. Fatigue strength also decreases with

increasing porosity [9]. Due to wide freezing range in AZ91 alloy, porosity is formed by lack of

interdendritic feeding in last stages of solidification [4]. Isothermal growth of eutectics effects in

decrease of alloy porosity [3].

It is important to investigate impact of the cooling rate and development of casting process to

achieve high quality castings, characterized by optimal microstructure.

Material for investigations

The material for investigations concerns 6 castings made from AZ91 magnesium alloy, fed and

cooled in different ways, cast gravitationally into sand mould. Chemical composition of the alloy is

presented in table 1.

Table 1. The chemical composition of the AZ91 alloy (wt. %).

Mg Al Zn Mn Cu Si Fe Ni Be

90.84 8.45 0.46 0.23 <0.001 <0.10 <0.003 <0.02 <0.0001


The arrangement of the casts in mould as well as the method of their feeding and cooling has been

presented in figure 1. The dimensions of the casts were equal 100x50x20mm, the samples for

examination have been cut out from the center of the cast that was cut into halves along the longer

axis of symmetry.

Fig.1 The scheme concerning the method of feeding and the cooling of the casts inside the moulds

and designation of samples.

The mould as well as the castings were made in the Zakład Metalurgiczny „WSK Rzeszów”. The

charge material consisted of 50% of the charge alloy AZ91 that was purchased at Magnesium

Elektron (MEL) and 50% of process scrap in form of feeders. The casting mould was made of self-

hardening, furan mass that consists of inhibitors such as sulfur, potassium fluoborate and boron

acid. The process of casting was conducted at the temperature 765 ± 5˚C and the mould was filled

in 12 seconds. The melt was performed in the atmosphere of protective gasses - Ar, CO2 and SF6.

The alloy was refined by: chemical refinement with the use of carbon compounds at the temperature

of 720˚C, the homogenization of the granule by way of overheating the alloy up to the temperature

of 860˚C, holding the alloy in that temperature for 10 minutes, fast cooling down to the casting

temperature - 765 ˚C.

Methodology of the research

The aim of the research was to examine the influence of feeding and the method of removal of the

heat in a mould on the microstructure as well as the metallurgical quality of the castings. Before

performing the microsections, radiographic analysis of the casting quality had been performed and

cooling curves from each of the casts had been registered. Macrophotographs that present the

distribution of the shrinkage porosity on the surface of the microsection were taken with an

stereoscopic microscope - Olympus SZX9. Qualitative analysis of the microstructure of the castings

was performed with the help of scanning microscope Hitachi S3400N, using the secondary

electrons (SE) and backscattered electrons (BSE) technique on unetched microsections and on light

microscope Olympus GX71 with the use of the bright field method, on microsections etched with

the reagent that consists of 10mm of HF and 90mm of H2O. The analysis of the chemical

composition was conducted by means of the energy dispersive X-Ray spectrometry (EDS) method

on microsections that did not undergo etching. Chemical composition of areas with higher levels of

porosity and areas without any defects were compared on the basis of the same samples.

Quantitative evaluation of metallographic parameters of the massive precipitates of the Mg17Al12

phase and the areas of appearance of this phase in plate-shaped and acicular form was performed on

microsections etched chemically with the reagent that consists of 10mm of HF and 90mm of H2O.

Quantitative evaluation of the grain size was performed on microsections that underwent chemical

etching with a reagent that consists of 20 ml of CH3COOH, 60 ml of C2H5OH, 1 ml of HNO3 and

19 ml of H2O. Decimal to binary conversion and the measurement of metallographic parameters

was conducted via a program used for image analysis called Met-Ilo v.12.1.

S1 - no cooler nor feeder

S2 - 20 mm thick cooler, no feeder

S3 - 20 mm thick cooler, with a feeder

S4 - 40 mm thick cooler, with a feeder

S5 - no cooler but with a feeder

S6 - 40 mm thick cooler, with no feeder


Results of the research

Casting defects. Increased porosity was detected on the sample surface of the following castings:

S1, S2 and S6. In fig. 2a) the porous surface (of S1) was compared with the flawless sample (S3).

Porosity in samples S1, S2 and S6 was characterized by heterogeneous arrangement, the shrinkage

pores occur in eutectic areas, inside of which congealed dendrites were visible (fig.2b)

Radiographic examination revealed the presence of casting defects also in the case of casting S5.

a)

b)

Fig. 2a) Sample S1 – visible microshrinkages (upper part), sample S3 – no microshrinkages (lower

part), b) porosity on the surface of sample S6, visible dendrites inside the voids, SEM, SE.

Areas of increased porosity were characterized by significantly lower concentration of aluminum

(table 2).

Table 2. The chemical composition of the flawless areas and the areas with increased shrinkage

porosity (wt. %)

Specimen Area Mg Al Zn

S1 microshrinkages 91.1 8.4 0.6

no microshrinkages 90.0 9.6 0.5





Microstructural analysis. The microstructure of the AZ91 alloy contains: solid solution of

aluminum in magnesium α, eutectics created by massive precipitates of Mg17Al12 intermetallic

phase and the solid solution α, clusters of lammelar and acicular precipitates of the Mg17Al12 phase

as well as fine precipitates of other intermetallic phases (fig. 3a). The analysis of the chemical

composition by means of the EDS method performed on these precipitates showed the presence of

particles rich in silicon and magnesium, manganese and aluminum as well as scarce particles that

contain rare earth elements. The observations performed with the use of light microscope on

microsections that underwent etching revealed strong microsegregation of the chemical composition

of the alloy with the alloying elements that segregate into eutectic regions as well as the precipitates

of the intermetallic phases (fig.3b). The regions where fine precipitates of the Mg17Al12 phase occur

were characterized by different degree of precipitate refinement. Precipitates with similar

refinement degree and morphology occurred in regions that were clearly separated from each other.

(fig.3c, d)

10 mm

10 mm


a)

b)

c)

d)

Fig. 3. The microstructure of the AZ91 alloy: a) precipitates of intermetallic phases, sample S5,

SEM, microsection that did not undergo etching, b) Chemical composition segregation, sample S6,

LM, microsection that underwent etching, c) sample S4, SEM, microsection that did not undergo

etching, d) the detail taken from figure 3c, SEM, microsection that did not undergo etching.

The structure of samples that did not have any casting defects and samples that exhibited

microshrinkage was different. In the first case massive precipitates of the Mg17Al12 phase indicated

a compact and discontinuous pattern (fig. 4a). The precipitates of this phase, in samples with

microshrinkages, were continuous and fine. Inside of them islands of α solid solution (fig. 4b) were

visible. Precipitates of the Mg17Al12 phase in lammelar and acicular forms, in the samples that did

not exhibit microshrinkages were significantly larger and more compact, separated by a visible

boundary from the α solid solution (fig. 4c). In the samples that had microshrinkages these

precipitates occurred only in the vicinity of eutectic regions, were smaller and less compact

(fig. 4d).

The structure of the alloy also revealed phases that contain silicon and magnesium - probably the

Mg2Si as well as Al8Mn5 that contains aluminum and manganese.

The analysis of the Al content in the α solid solution . The content of the aluminum contained in

the solid solution α ranges from about 4 weight % in the dendritic regions up to 10% in eutectic

regions (fig.5)


a)

b)

c)

d)

Fig. 4. The microstructure of the examined alloy, microsections that underwent etching, LM. a,c)

sample S3, b) sample S6, d) sample S2.

Weight %

Mg-K Al-K

S4(1)_pt1 95.0 5.0

S4(1)_pt2 93.1 6.9

S4(1)_pt3 91.9 8.1

S4(1)_pt4 89.8 10.2

Fig. 5. The chemical composition of the α solid solution in different micro-regions of the alloy.

Quantitative analysis of the microstructure. The metallographic parameters of the massive

precipitates of the Mg17Al12 phase were presented in table 3; the parameters of the lammelar and

acicular precipitates of the same phase were presented in table 4 while the results of the

measurement of the grain size in table 5. Gray columns constitute the results of the samples without

porosity.

Table 3. Metallographic parameters of the massive precipitates of the Mg17Al12 phase. Parameter symbol unit S1 S2 S3 S4 S5 S6

volume fraction VV [%] 5.49 6.31 3.40 3.46 4.67 6.06

relative area of boundary SV [µm2/µm

3] 0.36 0.66 0.37 0.44 0.29 0.72

variation coefficient VV ν (VV) [%] 34 15 37 33 44 20

variation coefficient SV ν(SV) [%] 88 84 98 99 102 84


Table 4. Metallographic parameters of the lammelar and acicular precipitates of the Mg12Al12 phase. Parameter symbol unit S1 S2 S3 S4 S5 S6

volume fraction VV [%] 21.70 16.80 26.10 33.00 20.20 14.40

relative area of boundary SV [µm2/µm

3] 0.14 0.24 0.11 0.09 0.10 0.25

variation coefficient VV ν (VV) [%] 24 18 31 19 45 21

variation coefficient SV ν (SV) [%] 82 81 91 90 87 79

Table 5. Metallographic parameters of the grain size. Parameter symbol unit S1 S2 S3 S4 S5 S6

grain size

area of flat section A [µm2] 2246 968 3403 1536 6037 1250

number of grain per

unit area NA [mm

-2] 264 606 189 351 104 447

relative area of grain

boundary SV [µm

2/µm

3] 0.051 0.085 0.045 0.061 0.032 0.076

heterogeneity of the grain size

variation coefficient A ν(A) % 117.0 117.0 72.1 120.0 87.8 97.3

variation coefficient

NA ν(NA) %

12.70 25.80 9.26 14.30 9.52 12.80

grain shape

shape factor ξ - 0.642 0.544 0.640 0.659 0.660 0.605

elongation factor δ - 1.71 1.75 1.57 1.61 1.65 1.64

heterogeneity of the grain shape

variation coefficient ξ ν(ξ) % 21.7 32.0 20.1 34.0 19.3 27.3

variation coefficient δ ν(δ) % 34.1 34.0 28.3 37.6 30.2 27.9

Solidification curves. The solidification curves designated for the liquidus-solidus range of

temperatures of each casting were presented in fig.6. The time in which the temperatures achieved

the liquidus-solidus level as well as the time in which each of the castings solidified was presented

in tab.6.

Fig. 6. The solidification curves in the temperature/time relation of each casting, the range of

liquidus-solidus temperatures


Table 6 The match-up of the time in which the temperatures reached the liquidus-solidus level for

each of the castings.

temperature time [s]

S1 S2 S3 S4 S5 S6

liquidus 13.5 11.5 10.0 23.5 26.0 8.5

solidus 206.5 61.5 209.5 203.5 634.0 45.5

solidification time 193.0 50.0 199.5 180.0 608.0 37.0

Summary Observation of the structure of the microsections that underwent etching showed a high level of non-uniformity of the α solid solution etching. The solid solution was more etched in the vicinity of eutectic regions and inclusions of the intermetallic phases. The examination of chemical content of the solid solution by the EDS method showed a diversification of the aluminum part from about 4 weight % in the dendritic regions up to 10 wt. % in the oversaturated solid solution that belonged to the eutectic regions. This indicates strong segregation of alloy elements into the eutectic fluid which solidifies at the end and creates eutectic mixture. The observation of the microstructure via the scanning microscope revealed shrinkage porosity that is characteristic for eutectic regions. Solidified dendrites were revealed inside the pores. This indicates the presence of dispersed shrinkage cavities that came into being in the last stage of crystallization in which the possibility of filling out the inter-dendritic regions by eutectic fluid had been blocked. The regions that are of higher porosity are characterized by aluminum depletion. The manner in which the heat is removed inside the mould is of significant importance in terms of the microstructure as well as the metallurgical quality of the casts. The analysis of the solidification curves indicated that the castings in which the cooling speed dropped in later stages of the solidification process (parabolic shape of the curve) did not indicate any casting defects, due to time extension in which the eutectic fluid could feed the interdendritic regions freely. The casts whose cooling speed was linear were characterized by microshrinkages. The castings that had casting defects exhibited a significantly higher volume fraction of massive precipitates of the Mg17Al12 phase and higher level of their continuity. The S5 sample, via the RTG examination, indicated discontinuities inside the cast and was also characterized by increased quantity of these precipitates in relation to the S3 and S4 castings. In case of the castings that had no defects the volume fraction of the areas, where the Mg17Al12 was present in form of fine precipitates, was higher than in the case of defective samples. The observation described above results from lower speed of cooling in castings and the time extension in which fine precipitates could come into being from the oversaturated solid solution. The conducted study indicates that high cooling speed, optimal from the perspective of mechanical features (fine grain), may contribute to unacceptable porosity.

Conclusions

1. The microstructure of the examined alloy contains a solid solution of aluminum in magnesium

α, eutectic mixture created by massive precipitates of the Mg17Al12 intermetallic phase, as well

as the α solid solution, fine precipitates of the Mg17Al12 phase and fine precipitates of the

intermetallic phases: Mg2Si and Al8Mn5.

2. The manner of heat removal inside the mould significantly impacts the microstructure of the

alloy. The casts that possess microshrinkages are characterized by more continuous and massive

precipitates of the Mg17Al12 phase. The occurrence of α solid solution islands, inside these

precipitates, was higher than in the case of samples without porosity. The casts whose

solidification curves were characterized by linear nature contained dispersed shrinkage cavities.

The casts with lower cooling speed in the last stage of the solidification process were flawless.

3. The reduction in the speed of the heat removal, in the last stage of the process of solidification,

influences the volume fraction of the areas of occurrence for fine precipitates of the Mg17Al12.

phase. This comes as a result of time extension in which these precipitates come into being from

the oversaturated solid solution which is formed due to the solidification of the residual fluid.


4. The speed of heat removal inside the mould affects the refinement of the grain – higher speed of

the heat removal makes the grain smaller.

5. The nature of the cavities and the depletion of the areas with increased porosity in relation to

aluminium indicate contractile porosity while no gaseous pores were detected.

6. The best technological version in terms of the quality of the casts is the application of the feeder

as well as the cooler that are 20 or 40mm thick. One should note that feeders as well as cooler

that are 40mm thick are characterized by the size of the grain, which is twice as small.

Acknowledgment

The present work was supported by the Polish Ministry of Science and Higher Education under the

research project No 6ZR7 2009C/07354.

References

[1] ASM Speciality Handbook.: Magnesium and magnesium alloys. ASM International, 1999

[2] T. Rzychoń, A. Kiełbus, J. Cwajna, J. Mizera, Microstructural stability and creep properties of

die casting Mg-4Al-4RE magnesium alloy, Materials Characterization 60 (2009) 1107-1113.

[3] A. K. Dahle, Y. C. Lee, M. D. Nave, P. L. Schaffer, D. H. StJohn, Development of the as-cast

microstructure in magnesium-aluminium alloys, Journal of Light Metals 1 (2001) 61-72

[4] Y. Wang, B. Sun, Q. Wang, Y. Zhu, W. Ding, An understanding of the hot tearing

mechanism in AZ91 magnesium alloy, Materials Letters 53 (2002) 35-39

[5] J. Adamiec, Assessment of high-temperature brittleness range of casted alloy AZ91, Materials

Science Forum 690 (2011) 41-44

[6] K. Meshinchi Asl, A. Tari, F. Khomamizadeh, The effect of different content of Al, RE and Si

element on the microstructure, mechanical and creep properties of Mg–Al alloys, Materials

Science and Engineering A 523 (2009) 1-6

[7] A. Kiełbus, T. Rzychoń, The intermetallic phases in sand casting magnesium alloys for

elevated temperature, Materials Science Forum 690 (2011) 214-217

[8] C. D. Lee, K. S. Shin, Effect of microporosity on the tensile properties of AZ91 magnesium

alloy, Acta Materialia 55 (2007) 4293-4303

[9] H. Mayer, M. Papakyriacou, B. Zettl, S. E. Stanzl-Tshegg, Influence of porosity on the fatigue

limit of die cast magnesium and aluminium alloys, International Journal of Fatigue 25 (2003)

245-256


Fractography and structural analysis of WE43 and Elektron 21 magnesium alloys with unmodified and modified grain size

Stanisław Roskosz1, a, Bartłomiej Dybowski1,b , Jan Cwajna1,c 1Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

[email protected], [email protected], [email protected]

Keywords: magnesium casting alloys, WE43 alloy, Elektron 21 alloy, microstructure, fractography, mechanical properties, quantitative metallography

Abstract. Magnesium alloys, thanks to their low density, are characterized by very high specific strength and specific stiffness. Due to acceptable mechanical properties, these alloys are widely used in automotive and aerospace industries for the elements such as: gearbox and engine housings, steering wheel columns or wheels. Because of a big responsibility of the elements made from Mg-RE alloys, it is important to investigate modification impact on properties of the magnesium alloys. The paper presents results of studies on properties of the WE43 and Elektron 21 casting magnesium alloys, modified in three different ways – according to Magnesium-Elektron (MEL) specification, 50% stronger modification and 100% stronger. For the comparison, unmodified alloys were also investigated. Investigations showed, that alloys modified according to MEL specification presents sufficient set of structural and mechanical properties. Further increase of amount of modifiers doesn’t let to significant increase of mechanical properties. Fractographic investigations showed many non-metallic inclusions on the fractures surface, which are result of faulty smelting process. Introduction

WE43 is an magnesium alloy with rare earth elements addition, which exhibit good connection of mechanical properties and creep and corrosion resistance [1]. High cost of the alloy, however, led to development of cheaper alloy – Elektron 21 – presenting similar properties up to 200˚C [2]. These materials are widely used in aeronautical and automotive industries, where the weight saving is essential [2, 3]. Strict design of the alloys chemical composition allows to obtain optimal microstructure and adequate properties of the material. Optimal combination of Nd and Gd, which are added to magnesium alloys in form of so-called “hardeners”[4], results in improved precipitation hardening response and reduced microshrinkage [2, 5]. Zr addition has effect in potent grain refinement, which leads to obtaining required mechanical properties [6, 7]. What is more, Zr does not form new phases in Mg alloys [6]. Due to high chemical reactivity of Mg and rare earth elements, smelting process is conducted in the atmosphere of shield gases [8]. Liquid alloy is refined with fluxes, which absorb MgO inclusions and form heavy compounds, that are sinking to the bottom of crucible. MgO inclusions in the castings destroy their continuity lowering mechanical properties and corrosion resistance. Application of filters is a widely used purification method, which prevents bigger impurities from getting into the casting [9, 10]. Due to high responsibility of elements made from these alloys it is important to investigate chemical composition impact on the material properties and development of strict smelting and casting processes, which result in high quality of the castings. Material for investigation

The material for investigation consisted of WE43 and Elektron 21 casting magnesium alloys, with their chemical compositions outlined in tab. 1. Both unmodified castings were made for every alloy and with the addition of modifiers: Zirmax, Gd-Hardener, Nd-Hardener and 14RE. The modified castings were performed in 3 variants of modification: according to the MEL specification, 50% more and 100% more of the modifier with regard to the prescriptions of Magnesium Elektron (MEL).


Table 1. The chemical composition of the investigated alloys (wt. %) alloy Mg Zn Zr RE Y

WE43 bal. 0.2-0.5 >0.4 2.4-4.4 3.7-4.3

Elektron 21 bal. 0.2-0.5 0.5 2.6-3.1 Nd 1.0-1.7 Gd

-

The castings were made in Zakład Metalurgiczny "WSK PZL Rzeszów" sp. z o.o. For every variant, samples for tensile strength tests according to the PN-91/H-88052 norm, as well as samples for linear shrinkage according to the DIN and TGL (103-2011) norms were gravity sand-casted.

Methodology

Both unmodified and modified alloys underwent static tension trial on Zwick 1474 static test machine and impact strength examinations on Charpy's impact machine in room temperature. Fractures that emerged during impact strength examinations underwent fractographic tests. The observations were made using Hitachi S3400N scanning electron microscope (SEM), using the secondary electrons (SE) technique. The examinations of chemical compositions of inclusions on the surface of the fractures were made using the energy dispersive X-Ray spectrometry (EDS) method. The structural tests were made on metallographic sections cut from castings made for linear shrinkage tests. The observations of the structure were conducted on Olympus GX71 light microscope using the bright field technique. The quantitative analysis of the eutectics was conducted on unetched sections, images for testing were registered with 500x magnification. The quantitative assessment of the size of the grain was conducted on etched sections in etchants of chemical composition presented in tab. 2. Decimal-to-binary conversion and measurement of the eutectics and grain size were made in Met-Ilo v. 12.1.

Table 2. Reagents used for etching alloy etchant chemical composition

WE43 15 ml HNO3, 85 ml H2O Elektron 21 14 g CrO3, 17.6 g HNO3, 100 ml H2O


The results of mechanical tests for each variant are presented in tab. 3. Fractographic tests revealed numerous impurities on the surface of fractures in most of the samples from Elektron 21 alloy and samples from WE43 alloy, modified according to the MEL Specifications. The impurities rich in alloying elements and oxygen were revealed (fig. 1a) and impurities that additionally included: Al, Si, Ti, S, K, Ca (fig. 1b). On the surface of the Elektron 21 alloy, modified + 100%, an additional precipitates rich in zirconium were noticed (fig. 1c). On the surface of fractures in unmodified Elektron 21 alloys and WE43 alloy (fig. 2a) first type inclusions exclusively were revealed, whereas second type impurities were present in Elektron 21 alloy specifically (fig. 2b). All the fractures, apart from the sample cast from the unmodified WE43 alloy showed secondary crackings (fig. 2c). Moreover, on the fractures in each variant of the Elektron 21 alloy and unmodified WE43 alloy the existence of voids was noticed (fig. 2d). The structural tests of every alloy showed the presence of increased porosity, localised in the eutectic areas (fig. 2e). In the WE43 alloy modified + 50% and Elektron 21 alloy modified + 100%, singular surfaces of cleavage were observed (fig. 2f).


Table 3. The mechanical properties of cast alloys, before and after modification. alloy Rm

[MPa] R0.2

[MPa] R0.2 / Rm

[-] A5

[%] U

[J/cm2] WE43 - unmod. 203 135 0.67 6.0 74.4 WE43 - mod. acc. to MEL 193 153 0.79 2.1 135.0 WE43 - mod. +50% 174 166 0.95 0.5 213.2 WE43 - mod. +100% 197 139 0.71 3.5 78.1 Elektron 21 - unmod. 177 99 0.56 6.1 102.6 Elektron 21 - mod. acc. to MEL 168 118 0.70 3.6 219.3 Elektron 21 - mod. +50% 157 126 0.80 2.3 127.5 Elektron 21 - mod. +100% 157 124 0.79 2.5 70.2

a)

b)

c)

Fig. 1. The EDS tests of inclusions revealed on the surface of fractures a) WE43 alloy, modified according to MEL, b and c) Elektron 21 alloy, modified + 100%

Pt1

Pt2 Pt3

Pt1

Pt2

Pt1


a)

b)

c)

d)

e)

f)

Fig. 2. a) The non-metallic inclusion on the surface of the fracture in WE43 alloy modified according to MEL, SEM, SE, b) impurities on the surface of the fracture in the Elektron 21 alloy, modified + 100%, SEM, SE, c) secondary crackings, WE43 alloy, modification + 100%, SEM, SE, d) voids in the fracture, WE43 alloy, unmodified, SEM, SE, e) porosity on the surface of the Elektron 21 alloy sample, modification +50%, LM, f) surfaces of cleavage, WE43 alloy, modification +50% SEM, SE.

The structure of the alloys consisted of the solid solution of alloying elements in magnesium α as well as eutectics created by Mg-RE inter-metallic phases and solid solution of Mg α. Inside the grains of modified alloys fine precipitates were revealed, probably the nuclei of crystallisation rich in zirconium. Additionally, in the structure of Elektron 21 alloy fine, acicular or lamellar precipitates were observed, usually to be found in the vicinity of eutectic areas. The tests revealed the porosity in inter-dendritic areas, in the structure of Elektron 21 alloy in every case, whereas in WE43 alloy a delicate porosity, only in the unmodified alloy sample was revealed. The grain in WE43 alloy, both before and after the modification, showed full characteristics of cellular morphology. In turn, the Elektron 21 alloy before the modification contained grains of both dendritic and cellular morphology and only grains of cellular morphology after the modification. The results of the quantitative analysis of the eutectics are presented in tab. 4. The measurements of size and shape of the grain in tab. 5 and 6.


Table 4. The results of quantitative evaluation of eutectic's metallographic parameters.

Alloy Summary

Surface fraction AA [%]


Linear fraction LL [%]

Relative area of grain boundaries SV [µm2/ µm3]

WE43 - unmod. 3.72 19.9 3.78 0.637 WE43 - mod. acc. to MEL 4.72 10.4 4.83 0.669 WE43 - mod. +50% 4.75 14.5 4.77 0.608 WE43 - mod. +100% 4.89 20.1 4.91 0.647 Elektron 21 - unmod. 2.73 18.7 2.54 1.005 Elektron 21 - mod. acc. to MEL 2.71 12.6 2.70 1.026 Elektron 21 - mod. +50% 2.42 15.7 2.46 1.020 Elektron 21 - mod. +100% 2.24 18.8 2.36 1.058

Table 5. The collation of the results of WE43 grain size measurements

Parameter symbol unit WE43 - unmod.

WE43 - mod. acc. to MEL

WE43 - mod. +50%

WE43 - mod.

+100% grain size


SV [µm2/µm3] 0.066 0.104 0.110 0.131


grain shape shape factor ξ - 0.553 0.647 0.666 0.680 elongation factor δ - 1.62 1.70 1.68 1.66

Table 6. The collation of the results of Elektron 21 grain size measurements

Parameter symbol unit E21 - unmod.

E21 - mod. acc. to MEL

E21 - mod. +50%

E21 - mod.

+100% grain size


SV [µm2/µm3] 0.045 0.118 0.120 0.113


shape factor elongation factor ξ - 0.642 0. 622 0. 617 0. 624 area of flat section δ - 1.65 1.67 1.71 1.66

The fractographic investigations revealed the existence of non-metallic inclusions on the surfaces of fractures. In unmodified alloys and WE43 alloy modified according to the MEL specifications, there exist inclusions containing alloying elements and oxygen, which constitute fragments of oxygenated nappe of liquid metal, included in the casting during flushing. In all the variants of the Elektron 21 alloys inclusions containing the following elements: Al, Si, S, K, Ca, Ti can be found. The source of these inclusions is in the fluxes used during smelting. Moreover, on the surface of the fracture in the Elektron 21 alloy modified +100% fine precipitates rich in zirconium, probably nuclei of crystallisation, were observed. This may bear witness to the change of the character of cracking from inter-crystallic to trans-crystallic with surfaces of cleavage. Porosity that was presented on the surface of fractures in the inter-dendritic spaces was created as a result of lack of feeding of those areas in the last phase of solidification.


The influence of modification on the structure of WE43 alloy is unambiguous. With the increase of the amount of the modifier the size of the grain decreases, which is caused by the increase of the quantity of added nuclei of crystallisation in form of Zirmax modifier. The increase of the amount of the modifier causes the increase of volume fraction of the eutectic areas. This is caused by the increase of the amount of alloying elements which with magnesium constitutes inter-metallic phases (fig. 3a). a) b)

c) d)

Fig. 3. a, b) Relation between the amount of the eutectics and the size of the grain and the modification variants of the WE43 and Elektron 21 alloys, c, d) the collation of the volume fraction of the eutectics with the heterogeneity of their distribution in the structure in the WE43 and Elektron 21 alloys. In case of the Elektron 21 alloy the influence of modification is ambiguous. The size of the grain decreases along with the increase of the amount of modifiers, apart from the +100 modification variant, with which the size of the grain insignificantly increases. In the case of the eutectics, along with the increase of the amount of the modifier, their volume fraction increases as well (fig. 3b). In both cases the addition of the modifier according to the MEL specifications cause significant break-up of the precipitates. With the greater volume fraction of the eutectics, the relative surface of the boundaries increased. The modification according to the MEL specifications significantly homogenised the distribution of the eutectics in the volume of the alloy, further increase of the amount of the modifier increased the heterogeneity of the structure even more (fig. 3c and d).

0

200

400

600

800

1000

1200

1400

1600

1800

0

1

2

3

4

5

6

WE43 -

unmod.

WE43 -

mod. acc.

to MEL

WE43 -

mod.

+50%

WE43 -

mod.

+100%

[µm

2]

[%]

eutectics volume fraction VV [%]

surface area of the grain plane section…

0

500

1000

1500

2000

2500

3000

3500

0

0,5

1

1,5

2

2,5

3

Elektron

21 -

unmod.

Elektron

21 -

mod.

acc. to

MEL

Elektron

21 -

mod.

+50%

Elektron

21 -

mod.

+100%

[µm

2]

[%]


surface area of the grain plane…

0

5

10

15

20

25

0

1

2

3

4

5

6

WE43 -

unmod.

WE43 -

mod. acc.

to MEL

WE43 -

mod.

+50%

WE43 -

mod.

+100%

[%]

[%]



02468101214161820

0

0,5

1

1,5

2

2,5

3

Elektron

21 -

unmod.

Elektron

21 - mod.

acc. to

MEL

Elektron

21 - mod.

+50%

Elektron

21 - mod.

+100%[%

]

[%]




Modification in each case increases the yield strength of the alloys until the +50% variant. In the case of the quantity of the modifiers of the 100% higher amount in relation to MEL, the yield strength decreases (fig. 4a and b). The investigations showed a significant decrease of tensile strength for MEL modification as well as +50% modification. The +100% variant caused its minor increase (fig. 4a and b). These results may, however, be compromised by the presence of impurities in the structure of the alloys. The worth of A5 elongation also showed analogous relation for both alloys (fig. 4c and d). The impact strength in the case of WE43 alloy was increasing until the +50% modification, after which it decreased, whereas for the Elektron 21 alloy in this respect, only the modification according to MEL was beneficial (fig. 4c and d). a) b)

c)

d)

Fig. 4. a, b) the dependency of the value of yield strength as well as tensile strength on the variant of the modification of the WE43 and Elektron 21 alloys, c, d) the relation between elongation and impact strength and the modification variant, WE43 and Elektron 21 alloys.

Conclusions

1. The modification of the alloys according to MEL causes the increase of the amount of the eutectic areas and their break-up and homogeneity. Further increase of the amount of the modifier causes minor increase of the volume fraction of the eutectics (WE43 alloy) or its decrease (Elektron 21 alloy)

2. The modification of the alloys according to MEL specifications cause significant grain refinement, further increase of the modifiers causes merely minor refinement of the grain size.

3. The modification of the alloys increases their yield strength until the +50% variant, +100% of the modifier with respect to the MEL specifications causes the decrease of this property.

020

40

60

80100

120

140

160180

155160165170175180185190195200205210

WE43 -

unmod.

WE43 -

mod.

acc. to

MEL

WE43 -

mod.

+50%

WE43 -

mod.

+100%

[MP

a]

[MP

a]

Rm [Mpa] Re [Mpa]

0

20

40

60

80

100

120

140

145

150

155

160

165

170

175

180

Elektron

21 -

unmod.

Elektron

21 - mod.

acc. to

MEL

Elektron

21 - mod.

+50%

Elektron

21 - mod.

+100%

[MP

a]

[MP

a]

Rm [Mpa] Re [Mpa]

0

50

100

150

200

250

0

1

2

3

4

5

6

7

WE43 -

unmod.

WE43 -

mod.

acc. to

MEL

WE43 -

mod.

+50%

WE43 -

mod.

+100%

[J/c

m2]

[%]

A5 U [J/cm2]

0

50

100

150

200

250

22,5

33,5

44,5

55,5

66,5

7

Elektron

21 -

unmod.

Elektron

21 - mod.

acc. to

MEL

Elektron

21 - mod.

+50%

Elektron

21 - mod.

+100%

[J/c

m2]

[%]

A5 U [J/cm2]


4. On the surface of the fractures impurities were seen in form of non-metallic inclusions. Their source is the oxygenated nappe of liquid metal and fluxes used.

5. Owing to the presence of the non-metallic inclusions, it is difficult to unequivocally define the impact of the modifications on the tensile strength of the tested alloys.

6. The modification according to the MEL specification ensures the correct microstructure of the WE43 and Elektron 21 alloys, as well as good mechanical properties. For reasons of economy, the usage of larger amounts of the modifiers with regards to the MEL recommendations is disadvantageous.

Acknowledgment

The present work was supported by the Polish Ministry of Science and Higher Education under the research project No 6ZR7 2009C/07354.

References

[1] A. Kiełbus, T. Rzychoń, Mechanical and creep properties of Mg-4Y-3RE and Mg-3Nd-1Gd magnesium alloy, Procedia Engineering 10 (2011) 1835-1840

[2] P. Lyon, I Syed, S. Haeney, Elektron 21 – An aerospace magnesium alloy for sand cast & investment cast applications, Magnesium, Edited by K.U. Kainer, WILEY-VCH (2007) 20-25

[3] B. Smola, I. Stulíková, J. Pelcová, N. Žaludová, Phase composition and creep behavior of Mg-Rare earth-Mn alloys with Zn addition, Magnesium, Edited by K.U. Kainer, WILEY-VCH (2007) 67-73

[4] N. Hort, Y. Huang, D. Fechner, M. Störmer, C. Blawert, F. Witte, C. Vogt, H. Drücker, R. Willumeit, K.U. Kainer, F. Feyerabend, Magnesium alloys as implant materials – Principles of property design for Mg-RE alloys, Acta Biomaterialia 6 (2010) 1714–1725

[5] J. Adamiec, Repairing the WE43 magnesium cast alloys, Solid State Phenomena, 176 (2011) 99-106

[6] M. Sun, G. Wua, W. Wang, W. Ding, Effect of Zr on the microstructure, mechanical properties and corrosion resistance of Mg-10Gd-3Y magnesium alloy, Materials Science and Engineering A 523 (2009) 145-151

[7] ASM Speciality Handbook: Magnesium and magnesium alloys. ASM International, 1999 [8] T. Rzychoń, A. Kiełbus, M. Serba, The influence of pouring temperature on the

microstructure and fluidity of Elektron 21 and WE54 magnesium alloys, Archives of Metallurgy and Materials, 55 (2010) 7-13

[9] W. Wang, G. Wu, M. Sun, Y. Huang, Q. Wang, W. Ding, Effect of flux containing YCl3 on the yttrium loss, mechanical and corrosion properties of Mg-10Gd-3Y-0.5Zr alloy, Materials Science and Engineering A 527 (2010) 1510-1515

[10] J. Wang, J. Zhou, W. Tong, Y. Yang, Effect of purification treatment on properties of Mg-Gd-Y-Zr alloy, Trans. Nonferrous Met. Soc. China 20 (2010) 1235-1239


Precipitate processes in Mg-5Al magnesium alloy

Andrzej Kiełbus1,a 1Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

[email protected],

Keywords: Mg-5Al magnesium alloy, annealing, precipitation, elevated temperature, microstructure.

Abstract. In this article, the impact of long-term annealing on transformation of microstructure, in

sand casting and die-casting Mg-5Al magnesium alloy was discussed. The Mg17Al12 phase of the

diverse morphology is the basic strengthening phase in Mg-5Al alloys. After sand casting

microstructure of Mg-5Al alloy consists of α-Mg solid solution with continuous and discontinuous

precipitates of Mg17Al12 phase. After die-casting, the structure is characterized by significant grain

refining of α-Mg solid solution, however Mg17Al12 phase, together with α-Mg solid solution, forms

fully divorced eutectic. The Mg17Al12 phase undergo decomposition and coagulation at the

temperature above 180°C and higher.

Introduction

Most of commercial magnesium alloys are based on Mg-Al binary equilibrium system. From

among those alloys, AM50 alloy and AZ91 alloy, which include from 5 wt % of aluminium to 9 wt

% of aluminium, have predominant significance in the aircraft and the automotive industry.

Especially, applying die-casting technology (HPDC) Mg-Al alloy is characterized by good

castability and reveals high level of mechanical properties [1]. For the sake of very good

mechanical properties at ambient temperature, and the low creep resistance it can be utilized only to

the temperature of 120° C [2, 3]. For this reason they find application in complicated and thin-

walled die-casting for the automotive industry and in the large overall dimension of gravity casting

for the aircraft industry [4].

Two types of Mg17Al12 precipitates occur in Mg-Al alloys: continuous and discontinuous. In most

cases, the precipitates occur simultaneously. The continuous precipitates are a result of nucleation

and growth of individual Mg17Al12 phase particles, which leads to changes in the matrix

composition. Whereas discontinuous precipitates nucleate on the boundaries of the solid solution

grains and when growing, they take the form resembling nodules [6]. Mg-Al alloys containing 5÷10

wt.-% of Al, are dominated by the continuous precipitations of Mg17Al12 phase. However, it has

been found that the morphology of precipitations of the Mg17Al12 phase in Mg-Al alloys depends on

the Al content [5] and temperature (Fig. 1). It has been shown that when [6]:

• at T<Tc1 temperature – only continuous precipitations of phase Mg17Al12 occur in the alloy;

• in the temperature range of Tc1<T<Td1 – both continuous and discontinuous precipitations of the

Mg17Al12 phase occur in the alloy, where as the temperature rises, the number of discontinuous

precipitations increases;

• in the temperature range of Td1<T<Td2 – only discontinuous precipitations of phase Mg17Al12

occur in the alloy;

• in the temperature range of Td2<T<Tc2 – again, both continuous and discontinuous precipitations

of the Mg17Al12 phase occur in the alloy – however at the same time, along with the rise in

temperature, the number of discontinuous precipitations increases;

• in the temperature range of Tc2<T<Ts (solubility limit temperature – solvus) – only continuous

precipitations of phase Mg17Al12 occur in the alloy.

Critical temperature Tc1 occurs mainly in alloys containing 18,8 at.-% of Al. Other temperatures

occur in all commercial alloys. In Mg-Al alloys, continuous precipitation is prevailing at a high

temperature (close to solvus line) and at a low temperature, whereas in the range of temperatures in-

between, discontinuous precipitation prevails [7].


Fig. 1. Influence of Al content and temperature on the morphology of Mg17Al12 phase [6].

Material and experimental methods

The material for the research was the casting Mg-5Al magnesium alloy. The chemical composition

of this alloy is presented in Table 1.

Table 1. Chemical composition of investigated magnesium alloy

Alloy Element, wt %

Al Sr RE Mn Mg

Mg-5Al 4,9 - - 0.45 balance

Sand casting was carried out at 700°C. Die casting was carried out using a hot-chamber machine.

The casting temperature was 650°C. Long-term annealing was conducted at three different

temperatures: 180°C and 250°C for 500÷5000h with air-cooling. An Olympus GX 71

metallographic microscope and a Hitachi S-3400N scanning electron microscope were used to study

the microstructure. Metallographic specimens were made in accordance with the methodology

developed in the Department of Materials Science [8]. Hardness tests have been performed with a

Vickers indenter.

Results

After sand casting microstructure of Mg-5Al alloy consists of α-Mg solid solution with precipitates

of two types of Mg17 Al12 phases (Fig.2). First one of massive morphology, together with solid

solution, forms partially divorced eutectic (continuous Mg17 Al12 + α-Mg) at the grain boundaries.

a) b)

Fig.2. Microstructure of Mg-5Al alloy after sand casting.

0 5 10 15 20Al (%at)

C

Tc1

Td1

Td2

Tc2

Ts

C+D

C+D

D

CT (°C)

α (Mg)

400

300

200

100

C - continuous precipitates of Mg Al

D - discontinuous

T , T - temperature of

continuous precipitationT , T -

17 12

c1 c2

d1 d2

precipitates of Mg Al

temperature of

discontinuous precipitation

17 12


The second one of plate morphology is created as a result of discontinuous diffusional

transformation (Fig.2a). Moreover, globular precipitates of Al8Mn5 phase occur in the alloy

(Fig.2b). After die-casting, the structure of Mg-Al alloy is characterized by significant grain

refining of α-Mg solid solution, however Mg17Al12 phase, together with α-Mg solid solution, forms

fully divorced eutectic on the grain boundaries of α-Mg solid solution (Fig.3).

Fig.3. Microstructure of Mg-5Al alloy after die-casting.

Long-term annealing (500h) of sand casting at the temperature of 180°C leads to decomposition of

plate structure, which is formed as a result of cellular growth and coagulation of precipitation of

Mg17 Al12 phase (Fig.4). In the zones of the increasing aluminium content, continuous precipitation

of Mg17 Al12 phase is started (Fig.4 and 5).

Fig.4. Coalesced precipitates of Mg17 Al12 phase

on grain boundaries in Mg-5Al alloy, after

annealing at 180°C/500h/air.

Fig. 5. The zones of continuous and coalesced

precipitates of Mg17 Al12 phase in Mg-5Al alloy,

after annealing at 180°C/500h/air.

The extension of annealing time to 5000h (Fig.6) or increasing temperature of annealing to 250°C

cause continuous growing and coagulation of precipitation of Mg17 Al12 phase (Fig. 7). The

precipitation processes during long-term annealing in die casting Mg-5Al magnesium proceed

similarly as in sand casting, but the difference is that discontinuous precipitation of Mg17 Al12 phase

is not observed in the first step of the process (Fig.8). After 4000h of annealing, a big, coagulated

precipitation of Mg17 Al12 phase makes the structure of alloy on the grain boundaries of α-Mg solid

solution (Fig.9).


Fig.6. Different morphology of Mg17Al12 phase

precipitates in Mg-5Al alloy after annealing at

180°C/5000h/air.

Fig. 7. Microstructure of sand casting Mg-5Al

alloy, after annealing at 250°C/5000h/air.

Fig. 8. Precipitates of continuous Mg17 Al12

phase in die-casting Mg-5Al alloy after

annealing at 180°C/500h/air.

Fig. 9. Coalesced continuous precipitates of

Mg17 Al12 phase in die-casting Mg-5Al alloy

after at 180°C/4000h/air.

Fig.10. The influence of temperature and time

of annealing on the hardness of the Mg-5Al

sand casting alloy.

Fig.11. The influence of temperature and time

of annealing on the hardness of the Mg-5Al die-

casting alloy.

The precipitation processes of the Mg17Al12 phase during long-term annealing of sand casts

results in the growth of the hardness of the alloy alongside with lengthening the annealing time

(Fig.10). However in die-casts the structure of the alloy undergoes degradation in the result of

precipitation and the coagulation of Mg17Al12 phase, what reduces hardness significantly (Fig.11).


Discussion

Sand casting Mg-5Al alloy is characterized by the structure of α-Mg solid solution with partially

divorced eutectic Mg17Al12 + α-Mg and continuous precipitates of Mg17Al12 phase in areas with

higher content of aluminium. Moreover, globular precipitates of Al8Mn5 phase occur in the alloy.

After die-casting, Mg17Al12 phase and α-Mg solid solution form fully divorced eutectic.

Precipitation of Mg17Al12 phase proceeds continuously and discontinuously. In sand casting, firstly,

as a result of discontinuous precipitation, precipitates of plate Mg17Al12 phase are formed. Process is

started on the grain boundaries of α-Mg solid solution and consist in cellular growth of plate

precipitation of Mg17Al12 phase in the direction of central part of the grain. The growth of the plate

precipitates is proceeding incessantly till the matrix of alloy gets the equilibrium composition.

Volume fraction of plate zones is growing together with extension time of annealing. The second

step is coagulation of plate precipitates and the beginnings of continuous precipitation of Mg17Al12

phase in zones of increasing content of aluminium. Further annealing causes growth and

coagulation of both types of precipitates (Fig.12).

Fig.12. Diagram of precipitation of Mg17Al12 phase in sand casting Mg-5Al alloy.

a) discontinuous precipitation of Mg17Al12 phase;

b) continuous precipitation of Mg17Al12 phase;

c) growth and coagulation of precipitates of Mg17Al12 phase.

Fig.13. Diagram of precipitation of Mg17Al12 phase in die-casting Mg-5Al alloy.

a) continuous precipitation and coagulation Mg17Al12 phase precipitates;

b) dissolving of small precipitates of Mg17Al12 phase;

c) growth and coagulation of precipitates of Mg17Al12 phase.

temperature, time

aluminium richareas

discontinuous precipitationof Mg Al phase17 12

a)continuous precipitation

of Mg Al phase17 12

coagulation ofMg Al phase17 12

b) growth and coagulation of Mg17 Al12 phase

c)

coagulated precipitates of Mg Al phase17 12

c)

temperature, time

coagulated precipitates of Mg Al phase17 12

small precipitates of Mg Al phase17 12

a) coagulation of Mg Al phase precipitates17 12

dissolving of Mg Al phase precipitates17 12

b)


However, in die-casting microstructure, decomposition is started with continuous precipitation of

Mg17Al12 phase, small precipitates dissolve in the matrix and the rest coagulate and consequently

are growing (Fig.13).

Conclusions

Results of the research can help in creating conclusions of experience feature, as follows:

1. The Mg-5Al alloy after sand casting is characterized by the structure of α-Mg solid solution with

partially divorced eutectic Mg17Al12 + α-Mg, continuous precipitates of Mg17Al12 phase and

globular precipitates of Al8Mn5 phase.

2. The Mg-5Al alloy after die-casting, is characterized by significant grain refining of α-Mg solid

solution with fully divorced eutectic on the grain boundaries of α-Mg solid solution.

3. The Mg17Al12 phase undergo decomposition and coagulation at the temperature above 180°C and

higher.

4. The precipitation and degradation processes of the Mg17Al12 phase during long-term annealing of

sand casts results in the growth of the hardness of the alloy alongside with lengthening the

annealing time. However in die-casts the coagulation of Mg17Al12 phase reduces hardness of the

alloy.

Acknowledgement



References

[1] A.K. Dahle, Y.C. Lee, M.D. Nave, P.L. Schaffer, D. StJohn, Development of the as-cast

microstructure in magnesium-aluminium alloys, J. of Light Met. 1 (2001) 61-72.

[2] A. A. Luo, Recent magnesium alloy development for automotive power train applications,

Mat. Sci. For. 419-422 (2003) 57-65.

[3] T. Rzychoń, A. Kiełbus, Microstructure and tensile properties of sand cast and die cast AE44

Magnesium Alloy, Arch. of Metall. and Mat. 53 (2008) 901-907.

[4] M. Avedesian, H. Baker, Magnesium and Magnesium Alloys (ASM Handbook 1999).

[5] D. Bradai, M. Kadi-hanifi, P. Zięba, W.M. Kuschke, W. Gus, The kinetics of the discontinuous

precipitation and dissolution in Mg-rich Al alloys, J. of Mat. Sci. 34 (1999) 5331-5336.

[6] M. Zhang, P.M. Kelly, Crystallography of Mg17Al12 precipitates in AZ91D alloy, Scr. Mat. 48

(2003) 647-652.

[7] D. Duly, J.P. Simon, Y. Brechet, On the competition between continuous and discontinuous

precipitations in binary Mg-Al alloys, Acta Meta. Mater. 43 (1995) 101-106.

[8] J. Adamiec, S. Roskosz, J. Cwajna, J. Paśko, Repeatable and reproducible methodology of the

AZ91 alloy structure evaluation, Arch. of Foundry Eng. 7 (2007) 95-100.


Influence of pouring temperature on castability and microstructure of QE22 and RZ5 magnesium casting alloys

Bartłomiej Dybowski1, a, Robert Jarosz 2,b, Andrzej Kiełbus1, c, Jan Cwajna1,d 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

2 ZM „WSK Rzeszów”, ul. Hetmańska 120, 35-078 Rzeszów, Poland

a [email protected],b [email protected], c [email protected], d [email protected]

Keywords: QE22 and RZ5 magnesium casting alloys, castability, simulation, microstructure, quantitative analysis, hardness

Abstract: This paper presents results of investigations on influence of pouring temperature on

castability and microstructure of QE22 and RZ5 magnesium alloys. In case of QE22 alloy, the

filling length of the liquid alloy increased with the increasing pouring temperature. In RZ5 no such

dependence was noted. This is probably caused by oxide films in the structure of material. Grain

refinement and eutectics volume fraction also didn’t present correlation with pouring temperature.

Introduction

Magnesium alloys, due to their low density (1.8 g/cm3) and high mechanical properties are widely

used in automotive and aerospace industries [1÷4]. Decrease of construction weight guarantee

decrease of fuel consumption [1,4]. What is more, magnesium alloys are characterized by good

technological properties, such as castability and weldability. The main problems of these alloys are

their unsatisfactory high temperature properties and poor corrosion resistance [1÷3]. This led to

development of magnesium alloys with rare earth elements and ziroconium additions, which can

work in temperatures up to 250˚C [1,2].The RZ5 alloy is characterized by very good castability and

weldability [2]. High content of zinc leads to strong solid solution hardening [5], which guarantee

acceptable mechanical properties of the alloy. Silver addition in QE22 alloy is increasing response

to the age hardening [1], which provides high mechanical properties up to 200˚C [6]. Due to

complexity of the magnesium castings, it is important to investigate influence of different factors on

magnesium alloys castability. Recent studies on magnesium alloys with Al addition revealed, that

the main factor influencing alloys fluidity is pouring temperature [7,8]. Unfortunately, high

reactivity of magnesium, increase of pouring temperature may result in formation of oxide films,

that can decrease alloy fluidity [9,10]. Next factors influencing on fluidity are: mould temperature,

chemical composition of the alloy and formation of intermetallic phases [11]. So far, there is lack of

investigations concerning fluidity of magnesium alloys with rare earth elements and zirconium

additions. The paper presents investigation results of influence of pouring temperature on fluidity

and microstructure of QE22 and RZ5 magnesium casting alloys.

Material for investigation

The material for the research were two unmodified casting magnesium alloys: QE22 and RZ5, with

chemical composition presented in Table 1. The alloys were gravity sand casted in ZM "WSK

Rzeszów". The schematic model of the casting is presented in Fig. 1. The spirals for castability test

of each alloy were poured in temperatures of 755˚C, 798˚C and 835˚C.

Table 1. Chemical composition of investigated magnesium alloys (% wt)

alloy Mg Zn Zr RE Ag Cu

QE22 bal. - 0.6 2.0 2.5 0.07

RZ5 bal. 3.5-5.0 0.4-1.0 0.8-1.7 - -


Fig.1. The diagram of the castability test spiral model.

Methodology

The simulation of pouring the mould and solidifying the castings for each alloy was conducted in

Magma5, using new calculation option - Solver 5. The physical, chemical and thermal properties of

the alloys were verified in Department of Materials Science within the Silesian University of

Technology. The results of the simulation were compered with the real castability of the alloys,

expressed via the length of cast spirals. From every casting samples the cut samples were mounted

in conductive resin. The preparation of microsections included: grinding using abrasive paper with

gradation 220, 320, 500 and 1200, polishing using diamond pastes with grain size 3 µm or 1 µm and

finishing, using the paste Al 2O3 with grain size 0.25 µm.

The microstructure testes were conducted on Olympus GX71 light microscope. The observations of

eutectic areas and casting defects were conducted on unetched microsections using the bright field

technique. The observations of the grain were performed on the etched microsections, in the reagent

containing: 4.2g picric acid, 10 ml H2O, 10 ml CH3COOH, 70 ml C2H5OH, using the technique of

polarised light. Detecting and measuring every ingredient of the microsrostructure was performed

on the program called Met-Ilo v. 12.1.

The alloy hardness measurements were conducted on Duramin A-300 tester, using the Vickers

method with 1N load. Five measurements were conducted for each sample. The measurements were

being performed on polished microsections.


The simulations of pouring and solidifying of metal in the form were conducted using the Solver 5

module, taking the parameters of Advanced Turbulence and Surface Tension into consideration.

The first parameter includes violent flow of liquid metal, that may occur in the runner system and

on the lining of the mould, the second one describes the surface tension between the liquid metal

and the mould which changes along with the temperature. These two parameters, together with the

viscosity of the alloy decreasing along with the temperature describe the effectiveness of the flow of

liquid metal. The Solver 5 module generates a special grid - "Mesh for Solver 5", which does not

cause local errors of simplifying the profile of runner system and the casting.

One significant parameter that has the influence on the results of the simulation is the effective

filling process that is meant to ensure the optimal pouring capacity. An improper choice of the

capacity may lead to the runner receptacle overflow, which is caused by the construction of the

calculation algorithm, without including the overflow.

The comparison of both simulated and real castability showed twice as high simulated castability

for the QE22 alloy than the real one, and for the RZ5 alloy about three times as high real castability

than the simulated one. In both cases the increase of temperature of pouring causes the increase of

simulated castability. In case of the castings, this tendency was true for the QE22 alloy specifically.

The results of castability testes are presented in Table 2.

Casting

Main

inlet

Inlet –Place of

input of the liquid

metal nappe –

needed in

simulation


Table 2. The results of simulated and real castability of QE22 and RZ5 alloys.

QE22

T [°C] 755 798 835

simulation

102 mm 132 mm 156 mm

casting

54 mm 55 mm 71 mm

RZ5

T [°C] 755 798 835

simulation

21 mm 23 mm 27 mm

casting

66 mm 71 mm 63 mm

The microstructure of both alloys consists of dendrites of α-Mg solid solution and the eutectics

created via precipitates of intermetallic phases as well as α-Mg solid solution. The finest grains in

the structure of both alloys were equiaxial. Figures 2a and 2b present the eutectics in both alloys,

while Fig. 2c and 2d - the grains of α-Mg solid solution. The results of the quantitative analysis of

the microstructure is presented in Tables 3 and 4.


a)

b)

c)

d)

Fig. 2. The structure of tested alloys:

a) the eutectics in the QE22 alloy, bright field; b) the eutectics in RZ5 alloy, bright field;

c) the grain in the QE22 alloy, polarised light; d) the grain in the RZ5 alloy, polarised light.

In the structure of both alloys, the presence of non-metallic inclusions was noticed, which probable

constitute fragments of oxygenated stream of liquid metal, included inside the castings (Fig. 3a, b

and c). In their proximity, the inductility of the material was observed (Fig. 3a) as well as

significantly increased inter-dendritic porosity (Fig. 3b). In the structure of alloys cast in 835˚C

temperature no non-metallic inclusions were observed. In the RZ5 alloy the porosity was observed

only in samples collected from the end of the spiral (Fig. 3d). in the QE22 alloy, the pores were

observed in the sample collected from the beginning of the spiral poured in the temperature of

798˚C as well.

Table 3. The basic stereological parameters of eutectics of both alloys.


Table 4. The basic stereological grain parameters of both alloys

a) b)

c) d)

Fig. 3. The casting defects in the structure of the alloys:

a) non-metallic inclusions, QE22, the end of the spiral poured in the temperature of 798˚C, LM;

b) non-metallic inclusions and increased porosity, the same sample, LM;

c) non-metallic inclusions, RZ5, the end of the spiral pored in the temperature of 755˚C, LM;

d) increased porosity, RZ5, the end of the spiral poured in the temperature of 835˚C, LM.


The results of the hardness measurements for every sample are presented in Fig. 4.

a)

b)

Fig. 4. The hardness of tested alloys.

The results analysis

The simulated length of the spiral strongly differs from the lengths of real castings. In the case of

QE22 alloy, the simulated spirals are significantly longer than the castings. This is probably cause

by the following factors: non-metallic inclusions distorting the flow of liquid alloy, the degree of

refinement and modification. It is impossible to include these parameters during the simulation. In

case of RZ5 alloy, the simulated length of the spiral is significantly lower than the length of the

castings. It appeared that the faulty calculation algorithm of the MAGMA program was the cause of

such a state. It is indeed surprising, since the non-metallic inclusions observed in the real castings

make it more difficult for the liquid metal to flow freely, which is indicated by microshrinkages and

discontinuity of the material in their direct proximity.

The impact of pouring temperature on the microstructure of the samples remains ambiguous. Along

with the increase of temperature, the size of the grain does not present monotonic change in any of

the cases (Table 4). The volume fraction of the eutectics does not change in an unambiguous way

either (Table 3). However, a significant increase in the grain size was noted in the case of samples

collected from the beginnings of the spirals (Figure 5a and b). This is caused by the close proximity

of the runner system, which considerably limits the speed of heat removal. The volume fraction of

the eutectics in the case of samples collected from the beginnings of the spiral is also increased (Fig.

5b and c). The lower speed of heat removal in these areas increases the segregation of alloy

elements into the remaining residual fluid time, thus increasing the number of eutectic areas.

The hardness of tested alloys does not present unambiguous dependency from the pouring

temperature. However, a certain dependency between the hardness and volume fraction of the

eutectics was noticed. In the case of QE22 alloy, the increase of volume fraction of the eutectics

causes the increase of the hardness of the alloy, whereby the R2 correlation coefficient equals only

0.42 (Fig. 6a). For the RZ5 alloy, the tendency proved to be the opposite - along with the increase

of the volume fraction of the eutectics, the hardness of the alloy decreases, the correlation

coefficient equals 0.76 (Fig. 6b).

50.60 49.22 47.98 46.04 44.55 48.48

0

10

20

30

40

50

60

755 798 835

Vic

ke

rs h

ard

ne

ss

T [˚C]

Vickers hardness - QE22 alloy

beginning end

54.28 53.78 52.06 48.54 63.1 49.26

0

10

20

30

40

50

60

70

755 798 835

Vic

ke

rs h

ard

ne

ss

T [˚C]

Vickers hardness - RZ5 alloy

beginning end


a)

b)

c)

d)

Fig. 5a) The dependency of grain size from the pouring temperature for the QE22 alloy;

b) the dependency of grain size from the pouring temperature for the RZ5 alloy;

c) The dependency of volume fraction of the eutectics from the pouring temperature, QE22 alloy;

d) The dependency of volume fraction of the eutectics from the pouring temperature, RZ5 alloy.

a)

b)

Fig.6. The dependency of the hardness of tested alloys from the volume fraction of the eutectics.

0

500

1000

1500

2000

2500

755 795 835

Are

a o

f g

rain

's f

lat

sect

ion

[μn

2]

T [˚C]

QE22 alloy

beginning end

0

500

1000

1500

2000

2500

3000

755 795 835

Are

a o

f g

rain

's f

lat

sect

ion

[μn

2]

T [˚C]

RZ5 alloy

beginning end

0

1

2

3

4

5

6

7

755 795 835eu

tect

ics

vo

lum

e f

ract

ion

[%]

T [˚C]

QE22 alloy

beginning

0

1

2

3

4

5

755 795 835eu

tect

ics

vo

lum

e f

ract

ion

[%

]

T [˚C]

RZ5 alloy

beginning end

R² = 0,4185

44

46

48

50

52

4 4,5 5 5,5 6

Vic

ke

rs h

ard

ne

ss

eutectics volume fraction [%]

QE22 alloy

R² = 0,7646

48

53

58

63

68

3,7 4,2 4,7

Vic

ke

rs h

ard

ne

ss

eutectics volume fraction[%]

RZ5 alloy


Conclusions

1. The influence of the pouring temperature on the simulated castability is unambiguous, along

with the increasing temperature, the castability increases as well. In the case of real castings,

a similar dependency showed only in QE22 alloy.

2. In the proximity of non-metallic inclusions in the structure, there are numerous

discontinuities of the material as well as increased porosity, which indicates more difficult

feeding of those areas into liquid metal.

3. The size of simulated castings was considerably different from the real castings. In the case

of QE22 alloy, the length of the simulated spiral was twice as high than in the real spiral,

which explains the existence of non-metallic inclusions in the structure of the alloy, which

blocks the free flow of the fluid.

4. In the case of RZ5 alloy, the simulated length of the spiral is three times as low than the real

one, which is cause by the faulty calculation algorithm.

5. The influence of the pouring temperature on the microstructure and hardness of the castings

is ambiguous. The volume fraction of the eutectics and grain size do not change

monotonically along with the increase of pouring temperature.

6. The grain size and volume fraction of the eutectics increase in the case of samples collected

from the beginnings of the spiral. This is caused by the close proximity of the runner system

which decreases the speed of heat removal in the form.

7. The hardness of the alloys depends on the amount of eutectic precipitates.

Acknowledgment



References

[1] ASM Speciality Handbook, Magnesium and magnesium alloys, ASM International, 1999.

[2] J. Adamiec, The assessment of impact of construction factors on weldability of MSRB alloy,

Materials Science Forum 690 (2011) 37-40.

[3] T. Rzychoń, A. Kiełbus, Microstructure and tensile properties of sand cast and die cast AE44

Magnesium Alloy, Archives of Metallurgy and Materials 53 (2008) 901-906.

[4] K. Meshinchi Asl, A. Tari, F. Khomamizadeh, The effect of different content of Al, RE and Si

element on the microstructure, mechanical and creep properties of Mg–Al alloys, Materials

Science and Engineering A 523 (2009) 1-6.

[5] B. Bronfin, A. Ben-Dov, J. Townsend, S. Mahmood, J. Vainola, S. Deveneyi, N. Moscovitch,

Advanced gravity casting magnesium alloys for the aircraft industry, Magnesium, Edited by

K.U. Kainer, WILEY-VCH (2007) 14-19.

[6] Magnesium-Elektron UK, Magnesium casting alloys datasheet: 440

[7] W. Qudong, L. Yizhen, Z. Xiaoqin, D. Wenjiang, Z. Yanping, L. Qinghua, L. Jie, Study on the

fluidity of AZ91_xRE magnesium Alloy, Materials Science and Engineering A271 (1999) 109–

115.

[8] Q. Hua, D. Gao, H. Zhang, Y. Zhang, Q. Zhai, Influence of alloy elements and pouring

temperature on the fluidity of cast magnesium alloy, Materials Science and Engineering A 444

(2007) 69–74.

[9] W. Wang, G. Wu, M. Sun, Y. Huang, Q. Wang, W. Ding, Effect of flux containing YCl3 on the

yttrium loss, mechanical and corrosion properties of Mg-10Gd-3Y-0.5Zr alloy, Materials

Science and Engineering A 527 (2010) 1510-1515.

[10] J. Wang, J. Zhou, W. Tong, Y. Yang, Effect of purification treatment on properties of Mg-Gd-

Y-Zr alloy, Trans. Nonferrous Met. Soc. China 20 (2010) 1235-1239.

[11] T. Rzychoń, A. Kiełbus, M. Serba, The influence of pouring temperature on the microstructure

and fluidity of Elektron 21 and WE54 magnesium alloys, Archives of Metallurgy and Materials

55, Issue 1 (2010) 7-13.


The influence of section thickness on microstructure of Elektron 21 and QE22 magnesium alloys

Michał Stopyra 1,a, Robert Jarosz 2,b, Andrzej Kiełbus 1,c

1Silesian University of Technology, 40-019 Katowice, Poland,

2 ZM „WSK Rzeszów”, ul. Hetmańska 120, 35-078 Rzeszów, Poland a [email protected], b [email protected], c [email protected]

Keywords: Magnesium alloys, Elektron 21, QE22, stepped casting test

Abstract. The paper presents analysis of section thickness’ influence on microstructure of Elektron 21 and QE22 magnesium alloys in the form of a stepped casting test. Solid solution grain size and volume fraction of eutectic areas were measured using light microscope and stereological methods. The results showed the significant increase of grain size caused by wall thickness and its slight decrease connected with the distance between analyzed section and the gating system. This relationship was confirmed using statistical methods. QE22 alloy demonstrated finer grain structure than Elektron 21 alloy as well as lesser susceptibility of grain size to solidification conditions

1. Introduction

The increasing requirements concerning fuel usage and fumes emission force the search for lighter and lighter construction materials. Magnesium alloys, due to their good mechanical properties, hold application in motor and aerial industry. Because their lattice is not easily deformable, casting alloys have the greatest application. The most important factors which limit the application of magnesium alloys are high chemical activity responsible for low resistance to corrosion and difficulties concerning technological process (melting, casting) as well as low creep resistance. The attempts to enhance creep resistance of Mg alloys have led to development of the alloys containing rare earth elements (RE). They form intermetallic phases, which are stable in elevated temperature and are characterized by good coherence with matrix, strengthen solid solution [1], which improve weldability and decrease hot cracking susceptibility through reduction of the range between liquidus and solidus temperatures [2]. These phases enable also age hardening, efficient especially in alloys containing Nd and Gd (cheaper neodymium reduces the range of solubility of expensive gadolinium and enables saturation at its lower content) [3]. The significant element in magnesium alloys which do not contain aluminium is Zr because of its impact on grain refinement [4]. The alloy with the greatest technical application from the Mg-RE-Zr group is Elektron 21 [5]. The QE22 alloy, apart from rare earth elements addition, includes silver. It functions similarly to RE – it increases mechanical properties in elevated temperature; however, at the same time it reduces corrosion resistance of the alloy [3,4]. According to the Hall-Petch equation, the yield strength depends in inverse proportion to the square root of the average size of the grain. On the other hand, because of slippage on the grain boundaries – which is one of the mechanisms of hot deformation – a small grain is not desirable in the alloys intended to work with in elevated temperatures. Thus, obtaining the optimal features requires a compromise and development of the technology enabling attainment of satisfactory features – recognition of the correlation between the conditions of crystallization and microstructure. In order to check influence of the section thickness on the size of the grain, stepped casting tests are conducted. The subject of this work is the analysis of results of such a test conducted for the Elektron 21 and QE22 alloys.


2. Material for the research.

The material for research comprised two sand casting magnesium alloys: Electron 21 and QE22. The chemical composition of both alloys is given in Table 1.

Table 1. The chemical composition of the tested alloys [wt. %] (acc. to MEL)

Alloy Zn Zr RE Ag Cu Mg

Elektron 21 0.2-0.5 saturated 2.6-3.1 Nd 1.0-1.7 Gd

- - balance

QE22 - 0.6 2.0 2.5 0.07 balance

3. Methodology of the research

Fig. 1 presents the dimensions of the model used in the experiment. Samples intended for research in the grain’s size and volume fraction of eutectics were drawn from the centre of each section; they were marked with the consecutive letters of the alphabet, starting with the section located closest to gating system (Fig. 2).

Fig. 1. Stepped casting model’s dimensions [mm].

Fig. 2. Stepped casting model with gating system and marked place of samples’ drawing.

The samples were prepared through grinding with sandpaper SiC of grit ranging from 320 to 1200. Then, they were polished on polishing wheel with diamond paste of the average grain’s size 3 and 1 µm. The pictures of structures were taken on light microscope Olympus GX71 in the bright field technique. The quantitative analysis of microstructure was made in Met-Ilo program. The measure of volume fraction of eutectics areas was made on unetched microsections (Fig. 3); good contrast between matrix and precipitates has allowed applying the automatic procedure. The measure of the grain’s size was made on the etched samples (Fig. 4). The composition of reagents is given in Table 2. The surface area of the grains’ cross-section not cut by edges of the picture was


measured. Because of the strong activity of the etchants attacking the grains’ boundaries and the presence of numerous artefacts derived from digestion process, the measurements were made semi-automatically and manually.

Table 2. The chemical composition of the used etchants Alloy Composition of the etchant

Elektron 21 14 g CrO3 + 17.6 g HNO3 + 100 ml H2O QE22 4.2g C6H3N3O7 + 10 ml H2O + 10 ml CH3COOH + 70 ml C2H5OH

a b

Fig. 3. Unetched microsections – Elektron 21 (a) and QE22 (b)

a b

Fig. 4. Etched microsections - Elektron 21 (a) and QE22 (b)

4. Results and discussion

Figure 5 presents the results of quantitative analysis of the grain’s size for particular sections of castings from both examined alloys and corresponding coefficients of variation (quotient of standard deviation and mean value expressed in percents). In compliance with expectations, the biggest grain was observed in the place where the thickness of the wall was the greatest (section E), and the smallest one at the end of the system (section G). Correlation between thickness of the wall and size of the grain is clearly seen in both cases for all sections, particularly in QE22 alloy – increase or decrease of thickness of the wall in consecutive sections is connected with, correspondingly, the increase or decrease of the grain’s size. In Elektron 21 alloy sections A and B do not fit this correlation; however, other factors have also influenced the grain’s size:

- chill under section G, which made it crystallise in the first instance and in the quickest pace; - feeder above section E; - position in relation to gating system.


It is worth noticing that QE22 alloy characterised itself by considerably more fine-grained structure than Elektron 21. The average area of the grain’s cut in QE22 alloy fitted into the range 771-1863 µm2 and in Elektron 21 alloy into 1615 – 5668 µm2. In both alloys the most heterogeneous structure was observed in the thickest section – E, that is proved by the highest values of variability magnitudes. Fig. 6 presents the results of quantitative valuation of volume fraction of eutectics in both alloys together with variability magnitudes. In QE22 alloy the volume fraction of eutectics decreases linearly together with the increase of distance from gating system, and after reaching the minimum in section E (5,6%) increases rapidly and reaches the maximum value in section G (11,4%). This tendency can be relevant to granularity of the structure, leading to extended area of the grains’ boundaries, and to segregation of composition related to the order of coagulation, which in turn was induced by the presence of the chill in the thinnest section (G) and of the feeder above the thickest section (E). In the Elektron 21 alloy the volume fraction of eutectic has also reached the maximum in section G; however, for the rest of the parts no simple correlation can be seen, which is probably connected with the smaller fraction of eutectics in the whole volume of the alloy. Excluding section G, the difference between the maximum value (D – 3,8%) and minimum one (C – 2,5 %) amounts to barely 1,3 %. Smaller differences between fractions in particular sections make the tendency more difficult to catch.

Fig. 5. Size of the grain for particular sections in Elektron 21 and QE22 alloys.

Fig. 6. Volume fraction of eutectic in Elektron 21 and QE22 alloys.

5. Statistic analysis of the results

In order to analyse the results a test for coefficients of correlation and multiple regression was conducted [6]. The influence of thickness of the wall and the distance from gating system were taken into consideration. While appointing the distance from gating system, the 0 value was assumed at the beginning of section A, and then the distance to the centres of consecutive sections was calculated. Data for calculation is presented in Table 3.

50

55

60

65

70

75

80

85

90

95

100

600

1600

2600

3600

4600

5600

6600

A B C D E F G

Variation c

oeffic

ient [%

]

Surf

ace a

rea [µ

m²]

Elektron 21

50

55

60

65

70

75

80

85

90

95

100

600

800

1000

1200

1400

1600

1800

2000

A B C D E F G

Va

ria

tio

n c

oe

ffic

ien

ti [%

]

Su

rfa

ce

are

a [µ

m²]

QE22

0

10

20

30

40

50

60

70

80

90

100

0

1

2

3

4

5

6

7

A B C D E F G

Variation c

oeffic

ient [%

]

Volu

me fra

ction [%

]

Elektron 21

0

10

20

30

40

50

60

70

80

90

100

0

2

4

6

8

10

12

A B C D E F G

Variation c

oeffic

ient [%

]

Volu

me fra

ction [%

]

QE22


Calculations start from creating inputs matrix X, consisting of values in the first two columns of table 3 and the additional column, in which all the values amount to 1. The second matrix Y comprises measured values of the surface area of grain’s cross-section for the particular alloy. The analysis of regression was made through appointment of vector of regression coefficients b (eq.1 and 2) for both alloys:

Table 3. Data for calculation.

Section thickness [mm] Distance from gating

system [mm]

Surface area of grain’s

cross-section [µµµµm2]

Elektron 21 QE22

10 16.6 5401 1325 5 49.45 4575 1202

10 81.95 4133 1350 15 114.75 4745 1417 45 146.9 5668 1863 10 175.6 3249 1008 5 198.1 1615 771

−== −

5052

8.16

7.71

YX)XX(b 21ET1T

21E (1)

−== −

4.1232

5.2

0.23

YX)XX(b 22QET1T

22QE (2)

In both cases very high values of multiple correlation coefficients were obtained, for Elektron 21 alloy Rw = 0.965, and for QE22 alloy Rw = 0.970. Verification of significance of the regression coefficients and correlation coefficient unfolded successfully on the level of significance α = 0.05. Thus, it can be assumed that in the examined castings there is statistically significant linear correlation between the size of the grain, thickness of the wall and distance from gating system, given in the equations (3) and (4)

GE21 = 5052 + 71.7T – 16.8D (3) GQE22 = 1232.4 + 23T – 2.5D (4)

where: GE21/GQE22 – the average surface area of the grain’s cross-section in Elektron 21/QE22 alloys [µm2] T – coefficient of the wall’s thickness [µm2/mm] D – coefficient of the distance from gating system [µm2/mm] As it could have been expected, thickness of the wall has greater impact on the grain’s size than the distance from gating system, and the direction of these changes is opposite. From the correlation given it stems that Elektron 21 is characterised by considerably bigger grain and greater susceptibility of the grain’s size to the condition of crystallisation. No similar correlation for volume fraction of eutectics has been concluded.


6. Conclusions

The size of the grain in both examined alloys increases linearly together with the increase of the wall’s thickness and decreases linearly with the increase of distance from the gating system. Impact of the wall’s thickness is stronger than impact of the distance. Elektron 21 is characterised by significantly bigger grain and greater susceptibility of the grain’s size to the analysed factors than QE22. The given correlations have been statistically confirmed. Volume fraction of eutectics in both alloys was greatest where thickness of the wall was the smallest. On the basis of the obtained results the correlation between volume fraction and thickness of the wall can not be stated. Various values in the consecutive sections of casting have been caused by other factors. The results obtained for QE22 alloy suggest that segregation of composition, induced by the order of coagulation, could have had considerable influence. In Elektron 21 alloy the differences between the measured values were too small and coincidental to speak about correlation. In the future the comparison of the obtained results and the simulated breakdown of temperature during crystallization are planned, as well as the analysis of the impact of heat treatment on the size of the grain in the examined castings.

Acknowledgement

The present work was supported by the Polish Ministry of Science and Higher Education under the research project No 6ZR7 2009C/07354

References:

[1] ASM Specialty Handbook: Magnesium and magnesium alloys. ASM International, 1999 [2] J. Adamiec J., Weldability of the MSRB Magnesium Alloy, Solid State Phenomena 176,

(2011) 107 – 118. [3] M.B. Kannan, W. Dietzel, C. Blawert, A. Atrens, P. Lyon, Stress corrosion cracking of rare-

earth containing magnesium alloysZE41, QE22 and Elektron 21 (EV31A) compared with AZ80, Materials Science and Engineering A 480, 2008.

[4] U.Kainer - Magnesium Alloys and Technologies. Wiley-VCH Verlag GmbH & Co. Kg aA, Weinheim, 2003.

[5] A. Kiełbus, T. Rzychoń, Structural stability of Mg–6Al–2Sr magnesium alloy, Solid State Phenomena (2011) 75-82.

[6] M.Maliński – Wybrane zagadnienia statystyki matematycznej w Excelu i pakiecie Statistica, Wydawnictwo Politechniki Śląskiej, Gliwice 2010 (in polish).


The influence of tin on the microstructure and creep properties of a Mg-5Al-3Ca-0.7Sr-0.2Mn magnesium alloy

Tomasz Rzychoń 1,a, Bartosz Chmiela1,b



Keywords: magnesium alloys, microstructure, Mg-Al-Ca-Sr alloys, creep properties

Abstract. The paper presents results of microstructural investigations and creep properties of Mg-

5Al-3Ca-0.7Sr-0.2Mn (ACJM53) and Mg-5Al-3Ca-0.8Sn-0.7Sr-0.2Mn (ACTJM531) magnesium

alloys in the as-cast condition. The microstructure of ACJM53 consists of α-Mg, (Mg,Al)2Ca - C36,

Al3Mg13(Sr,Ca), Al2Ca - C15 and AlxMny. Additionally, the CaMgSn phase is observed in the

ACTJM531 magnesium alloy. The addition of 0.8 wt% tin reduces the tensile strength at ambient

temperature and creep resistance at 180ºC.

Introduction

Magnesium alloys due to their low density are suitable materials for application in the automotive

and aircraft industries. Magnesium alloys based on the Mg-Al system have been studied extensively

for use in vehicles due to the weight savings they provide and also for their excellent castability

[1,2]. Commercial magnesium alloys, such as AZ91, AM50 and AM60 have limited application

because of poor creep resistance and poor mechanical properties at elevated temperature of 120ºC.

The cause of this phenomenon is a low-melting point Mg17Al12 phase, which is located at the grain

boundaries. Therefore, it is important to reduce the amount of Mg17Al12 phase and introduce

thermally stable precipitates at grain boundaries as well as in the grain interior by adding proper

alloying elements. It is well known that alloys of Mg-Al-Ca systems may provide significant

improvement in elevated temperature properties due to reduction of volume fraction of Mg17Al12

phase and the formation of Al-Ca and Mg-Ca intermetallic compounds [1-3]. The presence of

highly stable Laves phases at grain boundaries and in the interior grains has a positive effect on the

creep properties of Mg-Al-Ca alloys, however the phase composition of Mg-Al-Ca alloys after

solidification is still under discussion. In magnesium alloys where the Ca/Al mass ratio is smaller

than 0.8 only the Al2Ca phase is present, whereas for greater Ca/Al mass ratio than 0.8 both Mg2Ca

and Al2Ca phases exist in the microstructure [5,6]. The Al2Ca phase with an ordered cubic C15

structure is the most advisable intermetallic compound among the Laves compounds that occurr in

Mg-Al-Ca alloys due to their high structural stability. The high structural stability is connected with

the high density of states (DOS) per atom near the Fermi level. The Mg2Ca compound with a

hexagonal C14 type structure has a smaller structural stability in comparison to the C15 phase,

however, it is much higher than Mg17Al12 phase [7]. Luo et al. [8] found the existence of the

(Mg,Al)2Ca phase with an hexagonal structure in AMC503 alloy. Suzuki et al. [9] identified

(Mg,Al)2Ca phase as a new Laves phase with a C36 structure and argued that in the Mg-rich corner

of the ternary phase diagram, the Al2Ca phase cannot directly form from liquid during

solidification, but forms by transformation of the (Mg,Al)2Ca phase due to solid phase

transformation [3]. Recently, S.M. Liang et al. [5] reported that the Al2Ca phase can appear in the

structure of Mg-Al-Ca alloys. Strontium addition to the Mg-Al-Ca alloys improves the solid-

solution strength of the α-Mg phase by increasing the Al solute content and causes the formation of

the Mg17Sr2 phase, which is identified also as Al3Mg13Sr [10]. The addition of tin to magnesium

alloys containing aluminum and calcium causes the formation of CaMgSn phase, which may

improve the creep properties. However, this phase impairs mechanical properties at ambient

temperature [11]. In this paper, the effect of tin on the microstructure of Mg-5Al-3Ca-0.7Sr-0.2Mn

alloy and creep properties is presented.


Experimental Procedures

Two magnesium alloys with aluminum, calcium, strontium, tin and manganese were prepared and

their compositions, which were analyzed by X-ray fluorescence spectroscopy, are listed in Table 1.

Commercially-pure Mg, Al, Sn and Mn were used, strontium and calcium were added in the form

of Al-10 wt.% Sr and Al-85 wt.% Ca master alloys, respectively. Melting of the alloys was

conducted by induction melting in an Al2O3 crucible under the protection of an argon atmosphere.

The melt was held at 730°C for 3 min then poured into graphite moulds.

Table 1. The chemical composition of investigated magnesium alloys (wt %)

Al Ca Sr Mn Sn Mg

ACJM53 5.1 2.96 0.69 0.15 - Balance

ACTJM531 5.02 2.94 0.68 0.17 0.76 Balance

Microstructural observations of the alloys studied were carried out using optical microscopy,

scanning electron microscopy (SEM) and scanning transmission electron microscopy (STEM).

Microanalysis of intermetallic compounds were performed by using energy-dispersive X-ray

spectroscopy (EDS). The volume fraction of phases was measured by quantitative metallography.

The phase identification was performed by X-ray diffraction analysis (XRD). Constant load tensile

tests were performed at 180ºC and 60 MPa. Creep strain was measured by extensometers which

were attached directly to the gauge section of specimens. The length of the specimen was 100 mm,

the gage length was 60 mm and the diameter of the reduced section was 6 mm.

Results and Discussion

Microstructure. Figure 1 shows an optical micrograph of the as-cast ACJM53 magnesium alloy. It

can be seen that the microstructure of this alloy consists of solid-solution α-Mg and secondary

solidification compounds distributing at interdendritic areas. The interdendritic compounds show

two morphologies, bulky phase and irregular-shaped eutectic (Fig. 2). In addition, globular particles

inside the α-Mg grains are visible. Higher magnification STEM observation exhibits fine needle

shape particles distributing in the α-Mg grains as shown in Fig. 3. The length of these particles does

not exceed 0.5 μm.

Fig. 1. Optical micrographs of ACJM53 magnesium alloy.

The microstructure of the ACTJM531 magnesium alloy is similar to the microstructure of ACJM53

magnesium alloy and consists of bulky phase, irregularly-shaped eutectic in interdendritic regions

(Fig. 4) and needle-shape particles inside the α-Mg grains. Addition of 0.8 wt% tin causes the

formation of irregular, coarse precipitates, which are mainly located in the vicinity of the eutectic

(Fig. 5).


Fig. 2. Scanning electron microscope (SEM) images of ACJM53 magnesium alloy.

Fig. 3. Needle-shape precipitates inside the α-Mg grains in the ACJM53 magnesium alloy.

Fig. 4. Optical micrographs of ACTJM531 magnesium alloy.

Energy-dispersive analysis results showed that the irregular eutectic phase contains magnesium,

aluminum, calcium (Table 2). It should be noted that the magnesium content in the EDS analysis

may be overestimated due to the interaction between the electron beam and magnesium matrix.

Based on the X-ray diffraction analysis (Fig. 6) in combination with the EDS results, it can be

concluded that the irregular phase has a hexagonal crystal structure of Mg2Ca type (P63/mmc) and

the chemical formula of this phase can be written as the (Mg,Al)2Ca. It is well known that the

hexagonal (Mg,Al)2Ca phase in Mg-Al-Ca alloys has C36 or C14 Laves phase structures depending

on the Al/Ca ratio [3]. In this case, lattice parameters of (Mg,Al)2Ca compound (a0 = 5.83 Ǻ, c0 =

18.897 Ǻ) indicate the C36 structure [12].


Fig. 5. Scanning electron microscope (SEM) images of ACTJM531 magnesium alloy.

Table 2. Average chemical compositions of phases in the investigated alloys, at.%, EDS.

Phase Morphology Mg Al Ca Sr Mn Sn

α-Mg in ACJM53 Matrix 98.6 1.4

α-Mg in ACTJM531 Matrix 98.9 1.1

(Mg,Al)2Ca Irregular eutectic 69.0 20.5 10.2 0.3

Al3Mg13(Sr,Ca) Bulky phase 70.7 19.9 2.0 7.4

CaMgSn Irregular precipitates 76.0 5.2 6.7 1.4 10.8

AlxMny Globular particles 31.9 37.1 31.0

The bulky phase contains magnesium, aluminum, strontium and a minor amount of calcium (Table

2). The bulky compound was also observed in Mg-Al-Sr alloys and tentatively designated as

stoichiometry Al3Mg13Sr [13]. In previous investigations using the Rietveld method [14], it was

reported that Al3Mg13Sr can be isomorphous with Mg12Nd (tetragonal I4/mmm crystal structure).

Also, in this case, diffraction lines from a tetragonal compound with lattice parameters of a0 =

10.31 Ǻ, c0 = 5.93 Ǻ was observed in the X-ray diffraction pattern. Thus, it can be stated that the

bulky Al3Mg13(Sr,Ca) phase is located in the interdendritic areas.

The needle-shaped precipitates in the solid solution grains were not positively identified in this

investigation due to their small volume fraction. However, based on the research work in Ref [15],

it can be assumed that it is Al2Ca Laves phase with the C15 crystal structure.

In the ACTJM531 magnesium alloy, the CaMgSn phase with an orthorhombic crystal structure

(a0 = 7.86 Ǻ, b0 = 4.66 Ǻ, c0 = 8.74 Ǻ) was identified. The literature data and morphology of

CaMgSn phase suggests that these precipitates form as a primary solidification phase [16].

The content of aluminum dissolved in the magnesium matrix is about 1.4 at.% (Table 2).

According to Vegard’s rule, if the aluminum content in magnesium solid solution is 1.4 at.%, the

lattice parameters of α-Mg should be a0 = 3.2044 Ǻ, c0 = 5.2026 Ǻ [17] (for pure magnesium

lattice parameters are a0 = 3.209 Ǻ, c0 = 5.211 Ǻ). Meanwhile, on the basis of X-ray diffraction the

measured lattice parameters are a0 = 3.2074 Ǻ, c0 = 5.2081 Ǻ. These results may indicate the

dissolution minor amounts of calcium and strontium in the magnesium matrix, because the calcium

(197.4 pm) and strontium (215.1 pm) atoms have a larger atomic radius than aluminum atoms

(143.2 pm) and their presence in solid solution will increase the unit cell of magnesium. According

to the Mg-Ca and Mg-Sr phase equilibrium diagrams, the solubility of calcium and strontium in

magnesium is about 0.3 at.% and 0.1 at.%, respectively. Therefore, at such a low solubility in solid

solution these alloying elements were not detected during the SEM-EDS microanalysis.


Fig. 6. X-ray diffraction patterns of the alloys.

Table 3 shows the volume fraction of the intermetallic compounds in the tested alloys. With the

addition of tin, the volume fraction of the (Mg,Al)2Ca and Al3Mg13(Sr,Ca) phases was decreased as

a result of the formation of the ternary CaMgSn compound.

Table 3. Volume fraction of intermetallic phases (vol. %) and average grain diameter in the alloys

Alloy (Mg,Al)2Ca-C36 Al3Mg13(Sr,Ca) Al2Ca-C15 CaMgSn AlxMny

Average

grain

diameter

ACJM53 5.9 1.2 0.5 - 0.1 148 µm

ACTJM531 5.1 1.0 0.3 1.2 0.1 159 µm

Figure 7 shows grains of solid solution in the investigated alloys. It can be assumed that the tin

addition will cause grain refinement due to the primary solidification of the CaMgSn particles,

which may be a heterogeneous nucleus for magnesium grains. In the ACTJM531 magnesium alloy,

the average grain diameter is slightly greater than that of the ACJM53 alloy (Table 3). However

statistical analysis using the Kruskall-Wallis test showed that difference in the average grain

diameter is not significant. Therefore, addition of 0.8 wt% tin to the ACJM53 magnesium alloy

does not affect the solid solution grain size.

Creep properties. The creep tests in the present investigation were carried out at a temperature of

180°C and at a stress 60 MPa. The creep curves for the ACJM53 and ACTJM531 alloys are shown

in Fig. 8. It can be seen that the creep curves exhibit a well-defined primary stage and a secondary

stage. From the gradient of the secondary stage in the creep curves, the steady-state creep rate can

be calculated and the results are shown in Table 4. The addition of tin to the ACJM53 magnesium

alloy increases the steady-state creep rate and the creep strain. It could be expected that the tin

addition and the formation of the CaMgSn precipitates will improve the creep resistance of the

ACJM53 alloy due to the higher melting point of the CaMgSn compound in comparison to the

(Mg,Al)2Ca and Al3Mg13(Sr,Ca) phases. However, the CaMgSn phase is characterized by an

unfavorable morphology. After creep test at 180°C and 60 MPa, voids were observed in the


microstructure of the ACTJM531 magnesium alloy in the vicinity of the coarse CaMgSn

precipitates (Fig. 9). The presence of such voids will lead to rapid destruction of the material during

service.

a) b)

Fig. 7. Optical micrographs of ACJM53 alloy (a) and ACTJM531 alloy (b) with visible grains of

magnesium solid solution.

Fig. 8. Creep curves of as-cast ACJM53 and ACTJM531 alloys at 180°C and at 60MPa.

Table 4. Tensile and creep properties at 60 MPa and 180ºC of as-cast ACJM53 and ACTJM531

alloys.

Alloy Creep strain ε,

%

Steady-state creep

rate [1/s] UTS [MPa] YTS [MPa] El. [%]

ACJM53 0,33 5,7·10-10

135 104 2.1

ACTJM531 0,36 1,03·10-9

118 99 1.7


Fig. 9. Microstructure of the ACTJM531 magnesium alloy after creep testing at 180°C and 60 MPa.

Fig. 10. Microstructure of the ACJM53 magnesium alloy after creep testing at 180°C and 60 MPa.

The irregular eutectic (Mg,Al)2Ca and bulky Al3Mg13(Sr,Ca) phases seems to be stable during creep

at 180°C and 60 MPa, however microvoids localized in the interdendritic regions were also

observed. Moreover, during creep of the ACJM53 alloy, the precipitation of Al2Ca phase occurred

in the α-Mg solid solution (Fig. 10). It should be noted that the needle precipitates of Al2Ca were

also visible in as-cast state, but their length did not exceed 0.5 µm, whereas in samples after creep

tests, precipitates of the Al2Ca phase had a length of about 2 µm. Thus, the better creep resistance of

ACJM53 magnesium alloy, in comparison to the ACTJM531 alloy, is caused by a precipitation

process and the absence of coarse CaMgSn phase.

Conclusions

The microstructure of the ACJM53 magnesium alloy consists of the α-Mg solid solution, irregular

(Mg,Al)2Ca eutectic phase, bulky Al3Mg13(Sr,Ca) phases and needle-shape precipitates of Al2Ca

phase inside the α-Mg grains. Moreover, globular particles of the AlxMny phase are observed in the

matrix. The addition of 0.8 wt% tin reduces the volume fraction of the (Mg,Al)2Ca and

Al3Mg13(Sr,Ca) phases and causes the formation of primary CaMgSn particles. The coarse CaMgSn

phase adversely affects the creep resistance of the ACJM53 alloy. The higher creep resistance of the

ACJM53 magnesium alloy, in comparison to the ACTJM531 alloy, is caused by precipitation of

needle-shape Al2Ca phase and the absence of coarse CaMgSn phase.


Acknowledgments


strategical project No. POIG.01.01.02-00-015/09 (FSB-71/RM3/2010)

References

[1] H.-T. Son, J.-S. Lee, C.-S. Kang, J.-C. Bae, K. Yoshimi, K. Maruyama, The effects of yttrium

element on microstructure and mechanical properties of Mg-5 mass%Al-3 mass%Ca based

alloys fabricated by gravity casting and extrusion process, Mater. Trans. 49 (2008) 945-951.

[2] A.A. Luo, Recent magnesium alloy development for elevated temperature applications, Int.

Mater. Reviews 49(1) (2003) 13–30.

[3] A. Suzuki, N.D. Saddock, J.W. Jones, T.M. Pollock, Structure and transition of eutectic

(Mg,Al)2Ca Laves phase in a die-cast Mg–Al–Ca base alloy, Scr. Mater. 51 (2004) 1005-

1010.

[4] Roskosz S., Adamiec J., Błotnicki M.: Influence of delivery state quality on microstructure

and mechanical properties of as cast AZ91 Mg alloy. Arch. Foundry Eng. 7 (2007) 143-146.

[5] S.M. Liang, R.S. Chen, J.J. Blandin, M. Suery, E.H. Han, Thermal analysis and solidification

pathways of Mg–Al–Ca system alloys, Mater. Sci. and Eng. A 480 (2008) 365–372.

[6] R. Ninomiya, T. Ojiro, K. Kubota, Improved heat resistance of Mg-Al alloys by the Ca

addition, Acta Metall. Mater. 43 (1995) 669–674.

[7] D.W. Zhou, J.S. Liu, P. Peng, L. Chen, Y.J. Hu, A first-principles study on the structural

stability of Al2Ca Al4Ca and Mg2Ca phases, Mater. Lett. 62 (2008) 206.

[8] A.A. Luo, M.P. Balogh, B.R. Powell Creep and microstructure of magnesium-aluminum-

calcium based alloys, Metall. Mater. Trans. A 33A (2002) 567–574.

[9] A. Suzuki, N.D. Saddock, J.W. Jones, T.M. Pollock, Solidification paths and eutectic

intermetallic phases in Mg–Al–Ca ternary alloys, Acta Mater. 53 (2005) 2823–2834.

[10] A. Suzuki, N.D. Saddock, L. Riester, E. Lara-Curzio, J.W. Jones, T.M. Pollock, Effect of Sr

Additions on the Microstructure and Strength of a Mg-Al-Ca Ternary Alloy, Metall. Mater.

Trans. A 38A (2007) 420-427.

[11] B. H. Kim, S. M. Jo, Y. C. Leeb, Y. H. Park I. M. Park, Microstructure, tensile properties and

creep behavior of Mg–4Al–2Sn containing Ca alloy produced by different casting

technologies, Mater. Sci. Eng. A 535 (2012) 40– 47.

[12] S. Amerioun, SI Simak, U. Häussermann, Laves-phase structural changes in the system

CaAl2-xMgx, Inorg Chem. Mar 10;42(5) (2003) 1467-74.

[13] E. Baril, P. Labelle, M.O. Pekguleryuz Elevated Temperature Mg-Al-Sr, Creep Resistance,

Mechanical Properties and Microstructure, J. Metals 55 (2003) 34–39.

[14] T. Rzychoń, A. Kiełbus, G. Dercz, Structure refinement of the multi-phase Mg-Al-Sr alloy

Solid State Phenom. Vol. 163 (2010) 169-172.

[15] A. Suzuki, N.D. Saddock, J.R. Terbush, B.R. Powell, J.W. Jones, T.M. Pollock, Precipitation

strengthening of a Mg-Al-Ca–Based AXJ530 die-cast alloy, Metall. Mater. Trans. A 39A,

(2008) 696.

[16] C. Zhang, H. Cao, V. Firouzdor, S. Kou, Y. Austin Chang, Microstructure investigations of

directionally solidified Mg-rich alloys containing Al, Ca and Sn, Intermetallics 18 (2010)

1597-1602.

[17] D. Hardie, R.N. Parkins, Lattice Spacing Relationships in Magnesium Solid Solutions,

Philos. Mag. Vol. 4, (1959) s. 815.

[18] A. Kiełbus, T. Rzychoń, The intermetallic phases in sand casting magnesium alloys for

elevated temperature, Mater. Sci. Forum 690 (2011) 214-217.


On the oxidation behaviour of WE43

and MSR-B magnesium alloys in CO2 atmosphere

Roman Przeliorz 1,a, Jarosław Piątkowski 1,b 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland


Keywords: magnesium alloys, oxidation, microstructure, DSC

Abstract. The aim of the studies was to determine the oxidation kinetics of two magnesium alloys, i.e. WE43 and MSR-B, in CO2 atmosphere with and without the addition of 2 vol.% H2O. The rate of oxidation was measured by thermogravimetry in the temperature range of 530-580oC, i.e. below and above the eutectic melting point. The melting point of the eutectic mixture was determined by differential scanning calorimetry (DSC). The corrosion products were analysed by scanning electron microscopy (SEM) and X-ray microanalysis combined with EDS. Studies showed that on the WE43 alloy, a two-layer scale was formed, in which the outer part was composed of yttrium and magnesium oxides, while the inner part contained only yttrium oxide. The scale was found to preserve its good protective properties even above the eutectic temperature. Analysis of the results showed that on the MSR-B alloy, under a thin, uneven layer of scale, the process of internal oxidation occurred, and at a temperature of 580oC, the alloy underwent partial melting.

Introduction

Magnesium alloys are commonly used at room temperature, but they can also be used at elevated temperatures and in oxidising environments. At various stages of processing, such as heating the charge, casting, machining, recycling, etc., process conditions can cause undesired effects, which will change the chemical properties and deteriorate the state of the product surface layer [1-4]. In [5] it was observed that at the beginning, the reaction between magnesium and oxygen takes place at three stages: chemisorption of oxygen on the surface of magnesium, the formation and coalescence of oxide precipitates, the formation of compact oxide layer. In the environment of water vapour, magnesium undergoes corrosion, but the reaction is proceeding slowly. Studies on protection of magnesium alloys against ignition have been going on since early 50s of the last century [6]. Typically, magnesium alloys are melted under a protective layer of gases (CO2, SO2, and SF6) to prevent oxidation and burning up [7]. The disadvantage of this method is its harmful impact on the environment and the need to install complex equipment, which additionally increases the cost of products. Another solution is to increase the flash point and oxidation resistance of magnesium alloys by modification of the chemical composition. Beryllium and calcium appear to be effective elements improving the resistance to oxidation. It was found that 3-8 ppm beryllium can significantly increase the oxidation resistance of magnesium alloys [8]. An addition of Ca can increase the oxidation resistance of these alloys up to a temperature approaching the melting point. Fan et al [9] showed that the addition of 0.3 wt% Ca raises the flash point of pure magnesium by 120 K. However, despite a high flash point, magnesium alloys with the addition of beryllium and calcium have not found until now a practical use in industry, mainly due to low mechanical properties and toxicity of beryllium. Therefore, it is necessary to search for new alloys with high flash point. Recent studies have indicated that yttrium not only increases the oxidation resistance of magnesium alloys, but can also provide excellent mechanical properties. Yet, to ensure the formation of an effective outer layer of yttrium oxide Y2O3, the concentration of yttrium in the alloy should be greater than 8 wt.%, which again means higher production cost [10]. Recently, interest has been focused on the development of high-strength wrought and cast magnesium alloys, particularly alloys cast in semi-solid state (the thixocasting process) [2, 3].


The aim of this study was to compare the corrosion resistance of Mg-Y-RE and Mg-Ag-RE alloys in carbon dioxide and water vapour atmosphere at temperatures below and above the eutectic melting point.

Test materials and methods

Thermogravimetric studies were carried out on the WE43 and MSR-B magnesium alloys. The examined alloys differed in the content of yttrium and silver. The chemical composition of alloys is given in Table 1.

Table 1. Chemical composition of magnesium alloys.

Alloy Chemical composition, [wt.%]

Y Ag Zr RE Mg

WE43 4,3 - 0,3 3,4 rest MSR-B - 2,4 0,4 2,5 rest

RE- rare earth metals (Nd, Dy,Yb,Gd)

Samples with dimensions of 10x2x15 mm were polished with up to 1200 grit abrasive papers. Oxidation measurements were carried out using a thermobalance made by Setaram. The reaction atmosphere was gaseous CO2 mixed with 0.1% CO, dry and humidified. The gas was humidified with distilled water at 20°C PH2O=2⋅103 Pa. The gas flow rate was1.2 l/h. The investigations were carried out at 530 and 580oC, i.e. below and above the eutectic melting point. The characteristic transition temperature was determined by DSC, using high-temperature multi HTC calorimeter made by Setaram. Measurements were carried out in argon of 50N purity. The heating and cooling rate was 10oC/min. The reference substance was Al2O3. Thermodynamic calculations were performed using a HSC Chemistry Ver 4.1 computer programme [11]. The corrosion products (morphology and chemical composition) were examined with a HITACHI S-4200 scanning electron microscope coupled with Thermo Noran energy dispersive X-ray spectrometer, equipped with a SYSTEM SEVEN programme for microanalysis. The scale morphology was examined using secondary electrons signal images (SE).

Results

Calorimetric analysis of WE43 and MSR-B magnesium alloys

The calorimetric studies of WE43 and MSRB-B magnesium alloys showed that on the DSC curves there were two exothermic effects present during cooling (Fig. 1). The first effect obtained for WE43 alloy was responsible for the solidification of alloy matrix. The liquidus temperature was 634.3°C. The second thermal effect at 537.5°C corresponded to the eutectic transformation. The heat of transformation was -225.8 and -2.7 J/g, respectively (Fig. 1). For MSR-B alloy, the liquidus temperature was 638.6°C, and the temperature of eutectic transformation was 532.1°C. Compared to WE43 alloy, the amount of heat evolved during the solidification of alloy matrix was smaller, while the amount of heat necessary for the eutectic solidification was higher. The values of heat were -213.9 and -7.3 J/g, respectively (Fig. 1b).


Fig. 1. The DSC curves obtained during cooling of magnesium alloys: a) WE43, b) MSR-B.

Oxidation kinetics

The study of the oxidation kinetics of the WE43 and MSR-B magnesium alloys showed that the oxidation of WE43 alloy followed (approximately) a parabolic law (Fig. 2). The value of exponent "n" in a general equation for oxidation of metals and metal alloys:

tks

mp

n

⋅=

∆ (1)

a)

b)


where:

∆s

m– the unit gain in weight after time,

t, n – the exponent, kp – the exponential reaction rate constant, [g2·cm-4·s-1]

varied from 1.7 to 2.3 at 530 and 580oC, respectively. In the humidified gas, at a temperature of 530oC, with n=1,4, the run of oxidation curves has indicated that the WE43 alloy had good resistance to oxidation, even above the melting point of the eutectic mixture (Fig. 3).

0

0.2

0.4

0.6

0.8

1

0 2 4 6 8

Time, h

(∆(∆ (∆(∆m

/s),

*10

3 g

/cm

2

kl = 4,4x10-7

, g.cm-2

.s-1

kp = 1,7x10-11

, g2.cm

-4.s

-1

WE43, T=580oC

o MSR-B, T=530oC

MSR-B, T=580oC

x WE43, T=530oC

kp = 1,4x10-12

, g2.cm

-4.s

-1

kp = 5,5x10-13

g2.cm

-4.s

-1

Fig. 2. Oxidation curves for WE43 and MSR-B alloys in CO2 at temperatures of 530 and 580oC.

0

0.1

0.2

0.3

0.4

0 2 4 6 8

Time, h

(∆(∆ (∆(∆m

/s)x

10

3, g

/cm

2

kp = 1,5x10-12

, g2.cm

-4.s

-

1

kl = 8,4x10-9

, g.cm-2

.s-1

kp = 6,9x10-13

, g2.cm

-4.s

-1

x WE43, T=530oC

o MSR-B, T=530oC

Fig. 3. Oxidation curves for WE43 and MSR-B alloys in humidified CO2, Pap OH

31022

⋅= , at a temperature T = 530oC.


Detailed analysis of the oxidation curves plotted for the examined magnesium alloys has indicated that the oxidation behaviour of MSR-B alloy changed from the parabolic law to a linear dependence (Fig. 2). The linear course of the alloy oxidation occurred at 580oC, i.e. above the eutectic melting point. After the reaction, the sample was partially melted, assuming the shape of an elongated drop. In the environment of humidified gas at a temperature of 530oC, after the initial period of parabolic oxidation of the MSR-B alloy (n=2.3), the kinetics changed to linear (Fig. 3).

Scale morphology

Studies of scale morphology aimed mainly at the determination what effect the temperature and corrosive environment might have on its structure and porosity. This information was obtained examining the surface and cross-section of oxide layers under an electron microscope. The morphology of the scale formed on the examined magnesium alloys varied and depended on the chemical composition of alloy (Fig. 4). On the outer surface of the WE43 alloy, convex areas, cracks and flat spots with an evenly distributed corrosion occurred (Fig. 5). The thickness of the oxide layer was about 6 microns. The oxide layer was composed of globular and elongated grains. The inner part of the scale, with a thickness of about 2 microns, was compact and sticking firmly to the substrate (Fig. 6). The concentration of yttrium and magnesium in the outer layer was 95.5 and 4.5 wt.%, respectively, while in the inner layer it reached the values of 99.1 and 0.9 wt%, respectively. The morphology and chemical composition of the scale in humidified gas was similar to that obtained in dry gas. In the place where a crevice appeared in the inner layer, the alloy substrate was oxidised, while in the crevice an elevated concentration of magnesium was observed (Fig. 7).

a) b)

c) d)

Fig. 4. Outer surface morphologies of WE43 and MSR-B alloy samples after oxidation in CO2 dry and humidified,

T = 530oC: a) WE43/CO2, b) MSR-B/CO2, c) WE43, CO2/H2O, d) MSR-B, CO2/H2O.


a)

b)

c)

Fig. 5. Surface morphology of WE43 alloy after oxidation (a), (b) and (c) X-ray energy spectra (EDS) from selected areas marked in Fig. (a), CO2 atmosphere, T = 530oC.

a)

b)

c)

Fig. 6. Cross-section of the scale formed on WE43 alloy (a), (b) and (c) X-ray energy spectra (EDS)

from selected areas marked in Fig. (a), CO2 atmosphere, T = 530oC.


a)

b)

c)

Point Mg S Y

wt. % at. % wt.% at.% wt. % at. % 1 2,65 9,06 97,35 90,94 2 20,92 49,17 79,08 50,83 3 88,63 96,42 0,38 0,31 10,99 3,27 4 96,94 98,98 0,33 0,26 2,72 0,76 5 97,28 99,24 2,72 0,76

Fig. 7. Cross-section of the scale formed on WE43 alloy after oxidation in humidified gas a),

(b) and (c) X-ray energy spectra (EDS) from selected areas marked in Fig. (a), and (d) results of chemical analysis for areas marked in Fig. (a), T = 530oC.

The outer surface of MSR-B alloy was strongly folded (Fig. 8). Under the discontinuous layer of

scale, a region of internal oxidation was formed. The oxidised area had the form of bubbled fractures (Fig. 8).

a)

b)

Fig. 8. Surface morphology of MSR-B alloy after oxidation in CO2 (a), (b) X-ray energy spectrum

(EDS) from selected area marked in Fig. (a), T = 530oC.

d)


a)

d)

b)

c)

Point

Mg S Cl Ag Nd

wt.% at. % wt.%

at. %

wt.% at. % wt.% at. % wt.% at. %

1 55,49 86,36 22,26 7,81 22,25 5,83

2 83,93 94,26 1,89 1,61 1,05 0,81 13,12 3,32

3 98,17 99,58 1,83 0,42

4 54,93 86,21 20,95 7,41 24,12 6,38

Fig. 9. Cross-section of the MSR-B alloy after oxidation in CO2 (a), (b) and (c) X-ray energy

spectra (EDS) from selected areas marked in Fig. a), and d) results of chemical analysis for areas marked in Fig. (a), T = 530oC.

Discussion of results

The analysis of literature data [14, 15] indicates that on the surface of WE43 alloy the following reactions can take place:

,22 )(2 MgOOMgO g =+ 2530/,1,1029 moleOkJGo

Co−=∆ (2)

,3

2

3

432)(2 OYOY g =+ 2530

/,1113 moleOkJGo

Co−=∆ (3)

Additionally, the reaction of transformation of magnesium oxide into yttrium oxide is possible:

,2

1

2

3

2

332OYMgMgOY +=+ moleYkJGo

Co/,9,62

530−=∆ (4)

Under the experimental conditions, in the atmosphere of CO2 + 0.1% CO, the equilibrium partial

pressure of oxygen was pO2=10-17 Pa and thus was higher than the dissociation pressure of MgO and Y2O3 oxides. The pressure values as calculated by the HSC programme were 10-62 and 10-70 Pa, respectively, and hence reactions (2) and (3) were possible. According to the principle of thermodynamics, oxides of high dissociation pressure are formed at the oxidant-scale interface, while those of low dissociation pressure are formed in the inner part of the scale. Therefore, the oxide layer formed on WE43 alloy should be composed of MgO oxide in the outer part and of Y2O3 oxide in the inner part of the scale. A similar mechanism of the formation of oxide layers on magnesium alloys with yttrium has been presented in [12]. From the values of Gibbs free energy change it follows that initially magnesium and yttrium can oxidise at the same time. Gradually,


however, the MgO oxide should prevail in the scale due to the diffusion of Mg 2 + ions faster than the diffusion of Y3 + ions. The diffusion coefficient of Mg 2 + ions in the MgO lattice is DMg= 10-6 exp (-150.000/RT) [13], and of Y3 + ions in the Y2O3 oxide is DY = 10-9 exp (-300.000/RT) m2 / s [14]. Hence, at 530oC, DMg= 1.7⋅10-16, and DY= 3⋅10-26 m2/s. The selective oxidation of magnesium can promote the reduction of MgO by Y (reaction 4), because partial pressure of oxygen at the MgO/Mg interface is lower than the dissociation pressure of MgO, but higher than the dissociation pressure of Y2O3. According to [15], yttrium oxide has better protective properties than magnesium oxide. The MgO oxide is a semiconductor of "n" type with the interstitial cations of Mg 2 + [16]. Incorporating Y 3 + ions into the Mg 1 + x O oxide reduces the concentration of cationic defects and raises electron concentration in the conduction band. Thus, cation conductivity decreases and electron conductivity increases. As a result, the oxidation rate should be reduced. Introducing to the MgO lattice a more electropositive cation in the form of Y2O3 admixture can be described with the following reaction [17]:

232 5,0222 OOeYOY oMg ++′+= • (5)

232 32 OYMgOY Mgi +=+ ••• (6)

where: ••iMg - a doubly ionised atom of magnesium,

•MgY - a Y3+ ion in the lattice node normally occupied by Mg2+,

oO - an oxygen anion in the correct position,

e′ - an electron in the conduction band.

On the surface of MSR-B alloy under the conditions of isothermal oxidation, reaction (2) can proceed. Oxidation of silver under these conditions is not possible, because the dissociation pressure of Ag2O exceeds the partial pressure of oxygen in the atmosphere. The mixture of MgO and Nd2O3 oxides forms a discontinuous, porous layer (Fig. 9). Therefore, the dissolution of the oxidant in the alloy is possible, coupled with the formation of areas of internal oxidation. At a temperature of 580oC, catastrophic corrosion occurs. The presence of neodymium in the layer indicates the possibility of Nd2O3 oxide formation according to the reactions:

,3

2

3

432)(2 ONdONd g =+ 2530

/,4,1052 moleOkJGo

Co−=∆ (7)

and

,2

1

2

3

2

332ONdMgMgONd +=+ moleNdkJGo

Co/,5,17

530−=∆ (8)

In [18] it has been stated that an addition of neodymium can alter the oxidation kinetics from linear to parabolic, thus increasing the resistance to oxidation. Yet, the formation of a protective layer of Nd2O3 will be possible only when the neodymium concentration in alloy exceeds the equilibrium one. The equilibrium concentration of neodymium can be calculated from equation (7):

Nd

Mgo

C a

aRTG o

2/3

530ln−=∆ (9)

(for simplification, the Mg activity was replaced with a mole fraction).

At a temperature of 530oC, the concentration of neodymium reaches 6 wt.%. Since the concentration of neodymium in alloy has not exceeded 2 wt.%, initially magnesium was the first one to oxidise. If the concentration was higher than the equilibrium one, the neodymium oxide Nd2O3 would form. It has also been demonstrated in [17] that, despite the presence of diffusion barrier in the form of an Nd2O3/MgO oxide layer, its role in reducing the outward diffusion of Mg2+

cations is limited, especially at high temperature. The effect of water vapour probably consists in the formation of voids and pores in the scale, due to the formation of Mg(OH)2 hydroxide. The appearance of Mg(OH)2 in the outer part of oxide layer was observed in [18]. The immediate effect of the hydroxide presence can be abandoning the parabolic mode of oxidation process.


Conclusions

− The oxidation behaviour of WE43 alloy in dry and humidified carbon dioxide approached the parabolic law. This mode of reaction was prevailing near and above the melting point of eutectic.

− The oxidation mode of MSR-B alloy changed from the parabolic law to a linear dependence. At a temperature of 580oC, the sample was partially melted.

− The scale formed on WE43 alloy, consisting mainly of yttrium and magnesium oxides, was playing the role of a barrier to corrosion. The oxide layer on MSR-B alloy was composed of magnesium and neodymium oxides. In the layer, the process of internal oxidation occurred. The morphology of internally oxidised regions had the form of bubbled fractures.

References

[1] J. Medved, Ć. Primoz, Ć. Mrvar, M. Voncina: Oxidation Resistance of Cast Magnesium Alloys, Oxid Met (2009) 71, p. 257

[2] R. Lindström, L.G. Johansson, G.E. Thompson, P. Skeldon, J.E. Svensson: Corrosion of magnesium in humid air, Corrosion Science 46 (2004), p. 1141

[3] F. Czerwinski: The oxidation behaviour of an AZ91D magnesium alloy at high temperatures, Acta Materialia 50 (2002), p. 2639

[4] G. Baril, N. Pebere: The corrosion of pure magnesium in aerated and deaerated sodium sulphate solutions, Corrosion Science 43 (2001), p. 471

[5] T. Do, S.J. Splinter. C. Chen, N.S. McIntyre: The oxidation kinetics of Mg and surfaces studied by AES and XPS, Surface Science 387 (1997), p. 192

[6] J.F. Fan, G.C. Yang, S.I. Cheng, H. Xie, W.X. Hao, M. Wang, Y.H. Zhou: Surface Oxidation Behavior of Mg-Y-Ce alloys at high temperature, Met. and Materials Trans. 36 (2005), p. 235

[7] S.P. Cashion, N.J. Ricketts, P.C. Hayes: Cover gas protection for molten magnesium, J. Light Met. 2 (2002), p. 37

[8] Foerster G: U.S. Patent 4,543,234, (1985) [9] J.F. Fan, G.C. Yang, S.L. Cheng, H. Xie, W.X. Hao, M. Wang, Y.H. Zhou: Nonferrous Met.,

14, (2004), p. 1666 [10] N.V. Ravi Kumar, J.J. Blandin, M. Suery, E. Grosjean: Effect of alloying elements on the

ignition resistance of magnesium alloys, Scripta Mater. 49 (2003), p. 225 [11] HSC Chemistry Ver 4,1, computer programme Finland, (1998) [12] J.F. Fan, G.C. Yang, Y.H. Zhou,Y.H. Wei, B.S. XU: Selective Oxidation and the Third-

Element Effect on the Oxidation of Mg-Y Alloys at High Temperatures, Met. and Mater. Trans. 40A, (2009), p. 2184

[13] F. Czerwiński: The Oxidation Of Magnesium Alloys In Solid And Semisolid States, Metals & Materials Society, (2003), p. 30

[14] X.Q. Zeng, Q.D. Wang, Y.Z. Lu, W.J. Ding, C. Lu, Y.P. Zzu, C.Q. Zhai, X.P. Xu: Kinetic study on the surface oxidation of the molten Mg-9Al-0.5Zn-0.3Be alloy, J. Mater. Sc. 36 (2001) p. 2499

[15] X.M. Wang, X.Q. Zeng, G.S. Wu, S.S. Yao, L.B. Li: Surface oxidation behavior of MgNd alloys, Applied Surface Science 253 (2007), p. 9017

[16] R.J. Gaboriaud: Self-diffusion of yttrium in monocrystaline yttrium oxide: Y2O3, J. Sol. Chem. 35, (1980), p. 252

[17] R. Przeliorz, Oxidation of WE43 and MSR-B Magnesium Alloys in CO2 Atmosphere, Rudy i Metale Nieżelazne, 2011, p.138-145

[18] X.M. Wang, X.Q. Zeng, Y. Zhou, G.S. Wu, S.S. Yao, Y.J. Lai: Early oxidation behaviors of Mg-Y alloys at high temperatures, J. All. Comp. 460 (2008), p. 368


Galvanic corrosion test of magnesium alloys

after plastic forming

Joanna Przondziono 1,a, Witold Walke 2,b, Eugeniusz Hadasik 1,c


2 Silesian University of Technology, ul. Akademicka 2A, 44-100 Gliwice, Poland


Keywords: galvanic corrosion, magnesium alloys AZ61 and AZ80, extrusion.

Abstract. The purpose of this study was to evaluate resistance to galvanic corrosion of magnesium alloys AZ61 and AZ80. Resistance to galvanic corrosion was evaluated with additional application of aluminium alloy 2017A and 8Mn2Si steel as reference materials. The tests were carried out by means of potentiostat VoltaLab PGP 201 by Radiometer with application of Evans method. The tests were carried out in the solution with concentration of 0.01 M NaCl in ambient temperature. For comparison, the relations of the surface of magnesium alloys to aluminium alloys and steel (1:1, 5:1 and 10:1) was differentiated in the experiment. It was proved that AZ80 alloy features slightly higher corrosion resistance in contact with aluminium alloy and steel.

Introduction

Due to their physical characteristics and most of all high relative strength, magnesium and its alloys are used in aircraft and automotive industry. Due to intensive search for the best and most efficient production technologies by means of plastic forming of magnesium alloys, new possibilities of application of those materials are still showing up.

Development of magnesium alloys falls for the 90-ties of the XXth century. It is strictly related to decrease in vehicles weight. Magnesium is mostly used as a component of aluminium alloys, whereas magnesium alloys are used for production of pressure castings, the main customer of which is automotive industry [1-3]. Alloys after plastic forming feature higher mechanical properties in comparison to cast alloys, and their strength and formability can be formed through heat treatment, mainly age hardening. Despite unfavourable mechanical properties, application of alloys for plastic forming is modest and it makes only 1 % of annual magnesium production worldwide. The main problem connected with development of magnesium alloys processing by means of plastic forming is their limited plasticity.

Plastic forming of magnesium and its alloys can be carried out, depending on the content of alloy components, only in the restricted range of temperatures. Magnesium alloys, due to varied chemical composition, can undergo plastic forming at the temperature over 200°C. The reason for deformation of magnesium alloys in the elevated temperature is the increase in slip plane (at the temperature of 20°C there is only one slip plane). Magnesium alloys are mainly subject to extrusion forging and hot forging. Extrusion forging of magnesium alloys is carried out mostly in temperatures of 320÷450°C at the rate from 1 to 25 m/min. Recently, hydrostatic forging method has been under development, which will enable to carry out this process at lower temperatures and obtain higher grain size-reduction of magnesium alloys [3-7].

Application of magnesium alloys is limited to a great extent, due to low resistance to corrosion resulting from insufficient protection properties of oxide layer created on the surface in the oxidising atmosphere or the layer of hydroxides created in water solutions. Corrosion resistance of magnesium alloys mostly depends on the content of alloy components (e.g. aluminium, increased portion of which improves corrosion resistance) and impurities (e.g. iron and nickel, which then substantially decrease corrosion resistance) [8-12]. Due to high chemical activity of magnesium, it is subject to galvanic corrosion when contacts metals or alloys with higher corrosion potential


[13-16]. Galvanic corrosion of magnesium is the result of contact with various metals in corrosive environment and it creates huge limitations for application of magnesium alloys in automotive and aircraft industries. In theory, galvanic corrosion can be eliminated by insulating direct contact between the respective alloy and other conductive materials. Unfortunately, in industrial practice it is very difficult to obtain, especially when designing vehicles.

The quantity of corrosion devastation in galvanic cell not only depends on the position of metals in electrochemical series, but also on the size of the surfaces of contacting. Therefore, the purpose of this study was to evaluate the resistance to galvanic corrosion of magnesium alloys AZ61 and AZ80 after plastic forming contacting aluminium alloy 2017A and steel 8Mn2Si. The tests were carried out in solution with concentration of 0.01 M NaCl by means of potentiostat VoltaLab PGP 201 by Radiometer, with application of Evans method. The tests were carried out in ambient temperature. The relation of the surface of anode and cathode alloys was differentiated for comparison purposes in the experiment (1:1, 5:1 and 10:1).

Materials and testing methodology

Magnesium alloys AZ61 and AZ80 (samples: d = 14 mm and g = 1 mm every one of them) after plastic forming by means of extrusion forging were used in the tests. Reference material in galvanic corrosion tests was aluminium alloy 2017A used in aircraft and automotive industries and steel 8Mn2Si in wide application in production of connecting elements. Chemical composition of tested materials is presented in Tables 1-4.

Table 1. Chemical composition of magnesium alloy AZ61, % mas.

Zn Al Si Cu Mn Fe Mg 0.61 6.2 0.02 <0.01 0.22 0.002 reste

Table 2. Chemical composition of magnesium alloy AZ80, % mas.

Zn Al Si Cu Mn Fe Mg 0.34 8.2 0.02 <0.03 0.13 0.005 reste

Table 3 Chemical composition of aluminium alloy 2017A, % mas.

Cu Mg Mn Zn Si Fe Cr Al 4.2 0.66 0.43 0.12 0.41 <0.7 <0.1 reste

Table 4. Chemical composition of steel 8Mn2Si, % mas.

C Mn Si P S Al. N 0.08 1.86 0.73 0.014 0.010 0.022 0.007

Galvanic corrosion was evaluated with application of Evans method, and measurements were

carried out by means of potentiostat VoltaLab PGP 201 by Radiometer – Fig. 1.

Fig.1. Scheme of the galvanic corrosion test


The tests were carried out in solution with concentration of 0.01 M NaCl at the temperature T = 20 ±1˚C. In the experiment, the size of surface of magnesium alloys in relation to the size of surface of aluminium alloys and steel was differentiated (AZ61 – 2017A, AZ61 – 8Mn2Si, AZ80 – 2017A, AZ80 – 8Mn2Si) in the proportion: 1:1, 5:1 and 10:1. During the tests, magnesium alloys (AZ61 and AZ80) served as anodes, whereas aluminium alloy (2017A) or steel (8Mn2Si) served as cathode. Calomel electrode (SCE) served as the reference electrode. Current was increased until shorting potential was obtained (intersection point of voltagram curves). Shorting potential was marked as the system potential E, and current as system current I.

Test results review

Fig. 2 presents the relations between electrode potentials in the function of galvanic cell current: magnesium alloy AZ61 – aluminium alloy 2017A.

a)

b)

c)

Fig. 2. Change of electrode potentials of the alloy AZ61 and 2017A in the current function with surface relation a) 1:1, b) 5:1, c) 10:1


a)

b)

c)

Fig. 3. Change of electrode potentials of the alloy AZ80 and 2017A in the current function with surface relation a) 1:1, b) 5:1, c) 10:1

Next, the relation between the value of electrode potentials in the function of cell current: magnesium alloy AZ80 – aluminium alloy 2017 was presented in Fig. 3. Compilation of obtained values of potentials and currents is presented in Table 5.

On the basis of carried out test it was proved that connecting AZ61 alloy with 2017A alloy with the same surface amount creates a galvanic cell with higher value of current in comparison to connection of AZ80 alloy with 2017A alloy with proportionate surface. It is mainly caused by decreased content of Al, which directly influences corrosion resistance of magnesium alloys. Galvanic corrosion process intensity is dependent to a large extent on the difference of corrosion potentials present between metals creating the cell. The bigger the difference, the faster the corrosion process. Lower value of corrosion potential of AZ61 alloy in relations to AZ80 alloy is also responsible for higher corrosion resistance of AZ80 in contact with 2017A alloy.


Table 5. Results of galvanic corrosion resistance tests of the system: AZ61 alloy – 2017A alloy and AZ80 alloy – 2017A alloy

Galvanic cell Relation of

surfaces

Corrosion potential of the anode

EAZ61/EAZ80, mV

Corrosion potential of the cathode

EAl 2017A, mV

Corrosion potential of the system

E, mV

Corrosion current of the system

I, µA

AZ61-2017A 1:1 -1720 -703 -1427 273 AZ61-2017A 5:1 -1475 -980 -1443 180 AZ61-2017A 10:1 -1560 -906 -1461 101 AZ80-2017A 1:1 -1479 -1097 -1310 151 AZ80-2017A 5:1 -1490 -1088 -1346 98 AZ80-2017A 10:1 -1460 -950 -1428 80

Fig. 4. Change of the electrode potentials of AZ61 and 8Mn2Si alloy in the function of current with

surface relation: a) 1:1, b) 5:1, c) 10:1

Galvanic corrosion is also influenced by the amount of the surface of metals participating in corrosion process. The tests showed that with the increase of the surface of anode – magnesium alloy – one can observe substantial decrease of current value, and consequently decrease of


corrosion rate. Lower value obtained for AZ80 alloy proves its better resistance to corrosion in connection with 2017A alloy. Increase of the surface of alloy AZ61 and AZ80 (anode) by 10 times in relation to the surface of 2017A alloy (cathode) did not cause acceleration of corrosion rate of the anode alloy, which is presented in Fig. 2 and 3 and in Table 5.

For galvanic cells: magnesium alloy AZ61 – steel 8Mn2Si and magnesium alloy AZ80 – steel 8Mn2Si, current and voltage relations are presented in Fig. 4 and 5, respectively, and electrical values obtained on the basis of them, which describe corrosion, have been compiled in Table 6.

Fig. 5. Change of the electrode potentials of AZ80 and 8Mn2Si alloy in the function of current with

surface relation: a)1:1, b) 5:1, c) 10:1

For cells created between magnesium alloys and steel 8Mn2Si, similar relations as for connections of magnesium alloys and aluminium alloys can be observed. Also in this case connection between AZ61 alloy with steel 8Mn2Si with the same amount of surface caused creation of galvanic cell with higher value of current in relation to the connection of AZ80 alloy with steel 8Mn2Si with adequate surface. It must be stated, though, that the difference of potentials of anode and cathode is substantially bigger, which proves that corrosion processes will proceed faster. Corrosion current of cells created between the alloys of magnesium and steel is much higher, too. But also in this case magnesium alloy AZ80 features better resistance to galvanic corrosion, and increasing the surface of the anode will be a more favourable solution.


The tests also proved that an important factor influencing the course of galvanic corrosion is the rate of cathodic processes. It was proved that steel 8Mn2Si caused higher dissolution rate of AZ61 and AZ80 alloys than 2017A alloy – Tables 5 and 6.

Table 6. Results of corrosion resistance tests of the system: AZ61 alloy – steel 8Mn2Si and AZ80 alloy – steel 8Mn2Si

Galvanic cell Surface relation

Corrosion potential of the anode

EAZ61/EAZ80, mV

Corrosion potential of the cathode

E8Mn2Si, mV

Corrosion potential of the system

E, mV

Corrosion current of the system

I, µA

AZ61-8Mn2Si 1:1 -1370 -500 -1206 761 AZ61-8Mn2Si 5:1 -1400 -600 -1139 296 AZ61-8Mn2Si 10:1 -1500 -532 -1272 235 AZ80-8Mn2Si 1:1 -1508 -485 -1164 650 AZ80-8Mn2Si 5:1 -1545 -580 -1221 363 AZ80-8Mn2Si 10:1 -1570 -570 -1224 340

Summary

Due to low density, magnesium and its alloys are materials used mainly there, where the mass of construction or product is crucial. An essential problem for application of magnesium and its alloys in production of elements of machines and devices in various branches of industry is their low corrosion resistance. It is caused by high chemical activity which magnesium itself features. Elements made of magnesium alloys used as components of machines and devices often contact elements made of other types of material. Such connections are reasonable as far as application as such is concerned, but they also trigger a wide range of unfavourable physical phenomena.

The most frequent problem that occurs on the connections of two elements made of different materials is corrosion resistance. Due to low corrosion potential of magnesium in the galvanic series of metals, this element in such connections usually serves as anode, which causes its easy and relatively fast pulping. Therefore, it must be highlighted that one of the factors limiting application of magnesium alloys is their low resistance exactly to galvanic corrosion. In such connections with other metals, magnesium must be insulated. Whereas in situations when direct contact is unavoidable, and it happens with helical connectors, these elements should be separated by means of sufficiently big washers or filled with sealing mass in order to limit direct contact of those material to the greatest extent. Galvanic corrosion takes place in the atmosphere with relative air humidity over 60 %, whereas the rate and intensity of the process as such is dependent on the difference of corrosion potentials of connected metals. This relation is directly proportional, i.e. together with increase in corrosion potential difference of contacting metals with the same surface amount, the rate and intensity of corrosion processes increase, too. The relation of contacting materials surfaces also has a direct impact on the intensity of phenomena related to galvanic corrosion.

Performed tests enabled to draw the following conclusions: 1. The system: magnesium alloy – aluminium alloy, is a more favourable connection of two

metals than the system: magnesium alloy – steel, when resistance to galvanic corrosion is concerned. It is caused by lower difference of corrosion potentials, which takes place between anode and cathode for connection of magnesium alloy with aluminium alloy in comparison with the connection of magnesium alloy and steel.

2. From among magnesium alloys analysed in the study it is AZ80 alloy that features higher resistance to galvanic corrosion in contact with aluminium alloy or with steel. It is mainly connected with higher content of aluminium in its chemical composition. This element is used as alloy additive most of all to increase strength, improve castability and decrease shrinkage of magnesium alloys, but also to improve corrosion resistance.


3. Together with increase of anode surface (magnesium alloy), substantial decrease of corrosion current can be observed, and in consequence – decrease of corrosion rate. In industrial practice a situation when a small surface of cathode contacts a large surface of anode is more favourable.

Acknowledgements

Financial support of Structural Funds in the Operational Programme – Innovative Economy (IE OP) financed from the European Regional Development Fund - Project "Modern material technologies in aerospace industry", No POIG.0101.02-00-015/08 is gratefully acknowledged.

References

[1] L. Čížek, M. Greger, L.A. Dobrzański, I. Juřička, R. Kocich, L. Pawlica, Structure and properties of alloys of the Mg-Al-Zn system, Journal of Achievements in Materials and Manufacturing Engineering 32 (2009) 179-187.

[2] L.A. Dobrzański, T. Tański, L. Čížek, Influence of Al addition on structure of magnesium casting alloys, Journal of Achievements in Materials and Manufacturing Engineering 17 (2006) 221-224.

[3] A. Kiełbus, D. Kuc, T. Rzychoń, Magnesium alloys – microstructure, properties and application, Monograph, Modern metallic materials - presence and future, Department of Materials Engineering and Metallurgy, Katowice, 2009.

[4] E. Hadasik, Tests of metal plasticity, Monograph, Printing House of the Silesian University of Technology, Gliwice, 2008.

[5] R. Kawalla, Magnesium and magnesium alloys, Monograph, Metal processing, Plasticity and structure, Printing House of the Silesian University of Technology, Gliwice, 2006.

[6] M. Greger, R. Kocich, L. Čížek, Forging and rolling of magnesium alloy AZ61, Journal of Achievements in Materials and Manufacturing Engineering 20 (2007) 447-450.

[7] K. Bryła, J. Dutkiewicz, P. Malczewski, Grain refinement in AZ31 alloy processed by equal channel angular pressing, Archives of Materials Science and Engineering 40 (2009) 17-22.

[8] G.L. Maker, J. Kruger, Corrosion of Magnesium, International Material Review 38 (1993) 138-153.

[9] W. Walke, J. Przondziono, E. Hadasik, J. Szala, D. Kuc, Corrosion resistance of AZ31 alloy after plastic working in NaCl solutions, Journal of Achievements in Materials and Manufacturing Engineering, 45 (2011) 132-140.

[10] J. Przondziono, W. Walke, E. Hadasik, B. Jasiński, Electrochemical corrosion of magnesium alloy AZ31 in NaCl solutions, Acta Metallurgica Slovaca 16 (2010) 254-260.

[11] J. Przondziono, W. Walke, J. Szala, E. Hadasik, J. Wieczorek, Evaluation of corrosion resistance of casting magnesium alloy AZ31 in NaCl solutions, IOP Conf. Series: Materials Science and Engineering 22 (2011) 012017 1-12.

[12] J. Przondziono, W. Walke, A. Szuła, E. Hadasik, J. Szala, J. Wieczorek, Resistance to corrosion of magnesium alloy AZ31 after plastic working, Metalurgija (Metallurgy) 50 (2011) 239-243.

[13] J.X. Jia, A. Atrens, G. Song, T.H. Muster, Simulation of galvanic corrosion of magnesium coupled to a steel fastener in NaCl solution, Materials and Corrosion 56 (2005) 468-474.

[14] J.X. Jia, G. Song, A. Atrens, Experimental Measurement and Computer Simulation of Galvanic Corrosion of Magnesium Coupled to Steel, Advanced Engineering Materials 9 (2007) 65-74.

[15] G. Song, B. Johannesson, S. Hapugoda, D. StJohn, Galvanic corrosion of magnesium alloy AZ91D in contact with an aluminium alloy, steel and zinc, Corrosion Sci. 46 (2004) 955–977.

[16] J.X. Jia, G. Song, A. Atrens, Influence of geometry on galvanic corrosion of AZ91D coupled to steel, Corrosion Sci. 48 (2006) 2133–2153.


Creep resistance of WE43 magnesium alloy joints

Agata Kierzek 1, a, Janusz Adamiec 1,b 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

[email protected], [email protected]

Keywords: WE43, repair welding, heat treatment, creep test

Abstract. Magnesium alloys of Mg-Y-RE-Zr series are characterized by creep resistance up to a temperature of 250oC, and can work up to a temperature of 300oC. These properties allow for the application of alloys of Mg-Y-RE-Zr series for the elements of racing car engines operating in the conditions of high loads and temperatures. The requirement of high reliability components of aircraft propulsion system, with high strength and corrosion resistance, also led to the use of these alloys in the aerospace industry. Welding technologies in cast magnesium alloys are applied in order to repair defects in castings, occurring in the casting process, as well as to regenerate worn out castings. Joints made of magnesium alloys should have at least the same properties as a finished casting. The literature lacks information on the properties of joints welded of cast magnesium alloys.This work includes examination of influence of heat treatment on creep resistance of alloy WE43. Material for the study comprised joints made by the TIG method, welded in the cast state. Creep tests were carried out on joints without heat treatment and joints after heat treatment. The tests were performed at the temperatures of 200 oC and 250oC during 100h. It was found that there is an increase in creep resistance of the joints after heat treatment.

Introduction

Low density of magnesium alloys, along with their high specific strength and stiffness are the reasons why these materials are used in automotive and aerospace industries, allowing for lower fuel consumption by reducing the weight of the structure. One of the directions of cast magnesium alloys development, which will allow for their wider use in vehicles and aircraft, is to increase their resistance to creep [1]. The group of magnesium alloys with improved resistance to creep are the alloys of Mg-Y-RE-Zr series, which are characterized by resistance to creep up to a temperature of 250°C, and can work up to a temperature 300°C. These properties enable the application of alloys of Mg-Y-RE-Zr series in the elements of racing car engines operating in the conditions of high loads and temperatures. The requirement of high reliability components of aircraft propulsion system, with high strength and corrosion resistance, also led to the use of these alloys in the aerospace industry [2,3]. Welding technologies in cast magnesium alloys are applied in order to repair defects in castings, occurring in the casting process, as well as to regenerate worn out castings [4,5]. Welded joints should have at least the same properties as a finished casting, because only then the repaired casting will be able to work under the same conditions as the casting which does not require any repair. The literature lacks information on the properties of joints welded of cast magnesium alloys. This work includes examination of the influence of heat treatment on the creep resistance of joints welded of alloy WE43 (Mg-4Y-3RE-Zr) using the TIG method.

Research material

Joints of WE43 casting magnesium alloy were used for the research, with the chemical composition and properties presented in Table 1. The microstructure of the alloy in the state after casting (Fig. 1) consists of grains of α-Mg solid solution, precipitations of phases with various contents of alloy additions: Mg41Nd5, MgY, Mg24Y5, as well as precipitations of intermetallic phase β (Mg14Nd2Y) [6].


Fig.1. Microstructure of the WE43 alloy after casting

Table 1. Chemical composition and properties of the WE43 alloy

Chemical composition of WE43 alloy [% mas.]

melt Zn Si Cu Mn Fe Ni Li Zr Y Nd RE other

ASTM B80 - - - - - - - Min. 0,4 3,7-4,3 - 2,4-4,4 -

20091842 0,01 0,01 0,004 <0,01 0,002 0,004 0,01 0,51 3,7 2,2 0,96 <0,01

Mechanical properties [7]

Tensile strength [MPa] Yield strength [MPa] Elongation [%] HV3

230 178 7 85

Welded joints were made using the method of welding with nonconsumable tungsten electrode in argon shield (TIG). 10 mm thick test plates of WE43 alloy were butt welded in the state after casting. As an additional material, a wire of 2.4 mm diameter was used, with the chemical composition similar to the base material (tab.2). Welding parameters were summarized in Table 2.

Table 2. Technological parameters of welding of the WE43

WE

43

Welding current, [A] Arc voltage, [V] Linear energy of the arc, [kJ/cm]

120 14 3,0

Chemical composition of additional material [% mas.]

2009

3472

Zn Si Cu Mn Fe Ni Ag Li Zr Y Nd RE Inne

0,03 <0,01 <0,01 0,012 0,002 0,000 <0,01 <0,01 0,44 3,7 2,2 0,84 <0,01

The study was performed on joints welded in the state after casting as well as on welded joints heat treated after welding. Heat treatment was performed according to the manufacturer’s recommendations [7] and involved the processes of solution heat treatment 8h/525oC/air and ageing 16h/250oC/air (treatment T6). The material was welded in the state of delivery. The microstructures of of the weld area without heat treatment as well as after heat treatment are presented in Fig. 2.


Fig. 2. Research material: a) refinement of structure in a weld without heat treatment b) weld after heat treatment

Joint macrostructure consists of the native material, the heat-affected zone and the weld. Total melting was achieved, which allowed us to qualify the joints for further research. Grain fragmentation was found in the weld area of the joint without heat treatment. Solution heat treatments and ageing treatments resulted in the growth of grain in the weld, in comparison with the weld without heat treatment, as well as in the dissolution of the phases’ precipitations.

Research methodology and results

Creep tests were performed on Zwick Roell Kappa 50DS creep-testing machine located in the Department of Materials Science, Silesian University of Technology. Samples used in the tests were cut perpendicularly to the direction of welding. Weld area was located in the center of the axis of the sample with a diameter of 6mm and 70mm in length. The tests were carried out at the temperatures of 200oC and 250°C and a stress of 70MPa to 120MPa. The test duration was 110 hours. Measurement of deformation during the test was carried out continuously by means of extensometer. Table 3 shows the values of total deformation and minimum creep rate of the tested samples. Creep curves for welded joints made of alloy WE43 without heat treatment and after heat treatment are shown in Fig. 3 and 4.

Table 3. Results of creep tests of WE43alloy welded joints

Test parameters results

weld weld HT tempera-

ture stress

Strain after 100h, %

Creep speed, [s-1]

Strain after 100h, %

Creep speed, [s-1]

200oC 90MPa 0,418 9,18·10-10 0,35 5,49·10-10

120 MPa 0,874 2,85·10-9 0,399 1,21·10-9

250oC 70 MPa 1,248 2,31·10-8 0,437 6,66·10-9 90MPa 4,487 1,23*10-7 1,392 1,89·10-8

b) a)


Fig.3. Creep curves of WE43 alloy welded joint, research temperature: 200oC

(weld – weld without heat treatment, weld HT- weld after heat treatment)

Fig. 4. Creep curves of WE43 alloy welded joint, research temperature: 250oC

(weld – weld without heat treatment, weld HT- weld after heat treatment) The study of the weld area microstructure after a creep test was performed on Olympus GX9 light microscope in bright field technique [8]. Figure 5 shows the microstructure of the welds after the creep test.


Fig.5. Microstructure of weld after creep test (σ=90MPa,T =250oC) : a) cracks in weld without heat treatment b) weld after heat treatment

Results analysis and conclusions

The performed creep resistance tests allowed for the determination of creep curves for WE43 alloy joints. For all the creep curves obtained, one can distinguish a primary creep stage, characterized by a decrease in creep rate, and a secondary creep stage, in which creep rate is steady due to the balancing of the processes of consolidation and recovery. In the studied ranges of stress (70-120MPa) and temperature (200-250°C) after 110 h of creep test duration, there was no tertiary creep stage, characterized by rapid growth in the creep rate.

The minimum creep rate (Eq. 1), relative to the secondary stage of creep, which is responsible for the durability of the material, for metals and alloys is expressed in the exponential dependence [9,10]:

−=

RT

QAσε nexp

(1)

where: A - constant, σ - stress, n - constant, Q – activation energy, R – gas constant, T – temperature

Analysis of the results summarized in Table 3 indicates that heat treatment improves the creep resistance of joints welded of alloy WE43. The value of deformation after 110h of testing for heat-treated samples is even 3-fold lower than for thermally rough samples, tested at a temperature of 250°C. Deformation rate is also significantly lower for samples after solution heat treatment and ageing. Welded joints heat treated after welding are characterized by low value of deformation (ε = 0,39% at 200oC, ε =1,392% at 250oC), which allows for their use in the conditions of operative stress up to 120MPa at a temperature of 200°C and stress of not more than 90MPa at a temperature of 250oC. According to the data obtained from Magnesium Electron company [8] stress of 75MPa at a temperature of 250°C after 100 h of test duration, causes the deformation of 0.2% in the sample. In the case of the tested joints, stress of 70MPa causes deformation more than 2-fold higher (0.437%). This indicates a lower resistance to creep of WE43 alloy joints in comparison with the alloy WE43. Microstructure of the joint without heat treatment changed as a result of applying a temperature of 250°C for 100h, the precipitations of phases partially dissolved. The microstructure of the joint without heat treatment, tested at a temperature of 250°C and a stress of 90MPa, revealed numerous cracks in the weld spreading along the grain boundaries (Fig. 4), indicating a poor resistance to creep under these conditions. The heat treated joint, examined on the light microscope, revealed no

a) b)


significant changes in the structure of the weld, which indicates that heat treated joints are suitable for operating at that temperature. In effect of heat treatment, microstructure was changed. Fine dispersion phase Mg12NdY makes difficult movement of the dislocation and creep resistance is growth. Basing on the performed tests and the analysis of their results, the following conclusions were drawn: • Heat treatment of the joints of WE43 magnesium alloy improves their creep resistance.

Deformation of the heat treated joint, after 110h of testing, is several times smaller than the deformation of the joint without heat treatment. Creep rates are also lower in the case of samples after heat treatment.

• The observation of the microstructure of heat-treated joints on the light microscope revealed no changes in the structure. Studies should be extended to observations at higher magnifications in order to accurately describe the creep mechanism for welded joints.

• Castings of WE43 magnesium alloy, repaired by welding technologies and heat treated after welding, can be used at a temperature of up to 250oC and a stress of up to 90MPa; at lower temperatures (200°C) the joints carry the loads of up to 120MPa.

Acknowledgements

The study has been financed by the National Science Centre within the project No 2442/B/T02/2011/40 “Structure and properties of welded joints of cast magnesium alloys in simulated operating conditions

References

[1] B.L Mordike: Creep-resistant magnesium alloy, Mat. Sci. and Eng. A324 (2002) 103-112 [2] J.G. Wang, L.M. Hsiung et al.: Creep of a heat treated Mg–4Y–3RE alloy, Mat. Sci. and Eng

A315 (2001) 81–88, [3] M. Avedesian, H. Baker, Magnesium and Magnesium Alloys, ASM Speciality Handbook,

1999, [4] B. Ścibisz, J. Adamiec: Evaluation of susceptibility to hot cracking of WE43 magnesium alloy

welds in transvarestraint test, Arch. of metal. and mater. 55 (2010) 132-141, [5] A. Kierzek, J.Adamiec: Design factors influencing weldability of the Mg-4Y-3RE Cast

Magnesium Alloy, IOP Conf. Ser.: Mater. Sci. Eng. (2011) [6] T. Rzychoń, A. Kiełbus : Microstructure of WE43 casting magnesium alloy, J. of Ach. in

Mater. and Manuf. Eng. 21 (2007) 31-34, [7] Elektron WE-43, Data sheet 467, Magnesium Elektron, Great Britain, 2006 [8] A. Szczotok , S. Roskosz : New possibilities of light microscopy research resulting from digital

recording of images, Mat. Sci. 23 (2005) 559-565 [9] M.O. Pekguleryuz, A.A. Kaya : Creep resistant magnesium alloys for Powertrain applications,

Proceedings of the 6th International Conference Magnesium Alloys and Their Applications, Edited by K.U.Keiner, Weinheim 2004

[10] A. Kiełbus, T. Rzychoń : Mechanical and creep properties of Mg-4Y-3RE and Mg-3Nd-1Gd magnesium alloy, Proc. Eng. 10 (2011) 1835-1840


Impact of heat treatment on the structure and properties of the QE22 alloy welded joints

Agata Kierzek 1, a, Janusz Adamiec 1,b 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

[email protected], [email protected]

Keywords: magnesium alloys, welding, heat treatment, creep resistance

Abstract. The QE22 cast magnesium alloy containing silver, rare earth elements and zirconium is

characterized by high mechanical properties and creep resistance of up to 200°C. It is cast

gravitationally into sand moulds and permanent moulds. After the casting process, some defects can

be visible in the material, but they are repaired with use of overlay welding and other welding

techniques. The repaired cast should possess at least the same properties as the one which does not

require any repairs. The aim of this thesis was to determine the impact of the heat treatment on the

microstructure of the QE22 alloy welded joint. The creep resistance of the welded joints was also

analyzed.

Introduction

The magnesium alloys develop in four directions, which are associated with the reduction of the

weight of the elements, while maintaining or improving their existing properties. Most works are

devoted to research on magnesium alloys with rare earth elements, silver and strontium, which are

characterized by satisfactory creep resistance in temperatures above 250°C. These properties enable

one to minimize the weight and the moment of inertia of the structural elements of machines and

equipment operating in elevated temperatures [1,2]. Properties of the QE22 cast magnesium alloy

remain stable up to the temperature of 200°C, which allows one to apply this alloy in the aerospace

industry for the manufacture of bodies, engines or gearbox housings and in the automotive and

military industries [3,4]. Addition of silver in the QE22 alloy increases the strength properties and

the creep resistance. As solubility of silver in magnesium decreases when the temperature drops

(below 465°C), it is also possible to apply the precipitation hardening of the alloy [3]. The rare earth

elements are added in the form of dididium, i.e. a mixture containing 85% Nd and 15 % Pr, for the

purpose of increasing the creep resistance [1]. Zirconium modifies the size of the grain in the alloy,

increasing its mechanical properties in the ambient temperature [1].

The microstructure of the QE22 alloy in the post-casting state consists of a solid solution of

magnesium and a partially separated α-Mg + (Mg,Ag)12Nd eutectic mixture (fig. 1a). The heat

treatment improves the properties of the alloy in the room and elevated temperatures. It includes the

process of the solution heat treatment (8h/525oC/air) and ageing (16h/250

oC/air) [5]. Under the heat

treatment, the morphology of the precipitates of the (Mg,Ag)12Nd phase changes (fig.1b) [6].

Fig. 1. Microstructure of the QE22 alloy: a) structure after casting, b) structure after solution heat

treatment and ageing

a) b)


Welding and pad welding of magnesium alloys is applied as a method of repairing defects occurring

after the casting process: cracks, micro-shrinkages, etc. The defects are repaired with the methods

of manual welding, usually TIG. The most frequent reason for rejecting the repaired cast or

construction are the cracks occurring during the welding process. The QE22 alloy exhibits the best

weldability after casting [6,7]. An important factor determining the validity of application of

welding techniques for repairing or joining the cast magnesium alloys is the structural stability and

the stability of the properties of the joint in operating conditions. Its basic operating properties

include: creep resistance, resistance to fatigue in various conditions of load, and resistance to

chemical and electrochemical corrosion in elevated temperatures [1,8].

The aim of this paper was to investigate the effect of the heat treatment on the microstructure of the

QE22 cast welded joint and to determine impact of the heat treatment on the creep resistance of the

joints.

Materials used in the research

A welded joint of the QE22 cast magnesium alloy containing silver, rare earth elements and

zirconium (trade name: MSR-B [5]) was used in the research. The chemical composition and the

properties of the alloy were presented in table 1.

Table 1. Chemical composition and properties of magnesium alloys

Chemical composition of QE22 alloy [% of weight.]

Cast Zr RE Ag Other Mg

BS EN 1753 0.4-1.0 2.0-3.0 2.0-3.0 <0.01 rest

4377 0.46 2.57 2.4 <0.05 rest

Mechanical properties

Alloy Rm, MPa Re, MPa A5,% HV3

QE22 240 185 2 80

The welded joints were made with a welding method involving a nonconsumable electrode, in the

argon sheath (TIG). The test plates made of the QE22 alloy were butt welded in the post-casting state

(fig. 2). The additional material was a wire with of a diameter of 2.4 mm, whose chemical

composition was similar to the parent material (tab. 2). According to the manufacturer's

instructions, heat treatment was performed after the welding. The parameters of welding processes

and of the heat treatment are shown in table 2.

Fig.2. TIG method welding


Table 2. Parameters of TIG method welding

Welding parameters

Welding current, [A] Arc voltage, [V] Welding

speed,[cm/min]

Linear energy of the arc

[kJ/cm]

120 14 18 3.0

Additional material [% of weight]

Cast Ag RE Zn Mn Mg

20100666 2.4 2.3 0.01 0.02 rest

Heat treatment solution heat treatment: 8h/525oC/air + ageing 16h/250

oC/air

Methodology and results of the research

The performed joints were subject to visual examinations according to the PN-EN 970:1999

standard. They showed no welding inconsistencies in relation to both the face and the root of the

weld (fig.3a). Then the joints were subject to heat treatment. The samples used in the examination

of macro- and microstructure were cut perpendicularly to the welding direction, then they were

ground and polished with diamond pastes, according to the recommendations of the procedure

developed by the Department of Materials Science of the Silesian University of Technology [3].

The microsections prepared in this way were etched in 1% of nital. The macroscopic observations

were performed with the use of Olympus SZX9 stereoscopic microscope at magnification of 20 x

with the dark field technique. It was found that the butt joint had been made correctly with full joint

penetration (fig. 3b). The microstructure observations were made with the use of Olympus GX71

optical microscope with the bright field technique. Figure 4 presents the microstructure of non heat-

treated joint and the heat- treated joint after welding.

Fig.3. QE22 alloy joint a) weld face, b) macrostructure of joint

a) b)

a) b)


Fig.4. QE22 alloy welded joint: a) heat-affected zone – non heat- treated joint, b) non heat-treated

weld, c) heat-affected zone after heat treatment, d) heat- treated weld

The creep tests were performed with the use of the Zwick Roell Kappa 50 DS creep-testing

machine. The samples cut perpendicularly to the welding direction were used in the test. The weld

area was located in the centre of the axis of the sample with a diameter of 6 mm. The tests were

performed at the temperature of 180 °C and 200 °C and stress from 60 MPa to 90 MPa. The test

time was 110 hours (tab. 3). The strain was measured during the test with the use of an

extensometer. The creep curves for the QE22 alloy welded joints in the state without heat treatment

and after heat treatment were shown in figures 5. Table 3 presents the values of total strains and

minimum creep rate of the tested samples. Figure 6 shows the microstructure of the weld area after

the creep resistance test.

Fig 5. Creep curves for the QE22 alloy welded joints: a) examined at the temperature of 180°C,

b) examined at the temperature of 200°C, the lines 60MPa_weld and 70MPa_weld overlap,

(weld – non heat-treated joint, weld HT – joint after heat treatment)

Table 3. Parameters and results of the creep test of the QE22 alloy welded joints

Test parameters Results

Weld Weld HT

Temperature Stress Strain after

100h, %

Creep speed,

[s-1

]

Strain after

100h, %

Creep speed,

[s-1

]

180oC

70MPa 0.55 5.37* 10-9

0.19 2.76* 10-11

90 MPa 1.79 3.12* 10-8

0.31 9.48* 10-10

200oC

60 MPa 2.46 4.63* 10-8

0.21 9.72* 10-10

70MPa 2.46 4.56* 10-8

0.26 1.55* 10-9

c) d)

a) b)


Fig. 6. Heat-affected zone of the QE22 alloy welded joint after creep resistance tests: 60MPa,

200oC: a) state after casting, b) state after solution heat treatment and ageing

Analysis of results and conclusions

The QE22 alloy welded joint in the non heat-treated state after welding is characterized by correct

shape of the heat-affected zone. One can see a clear boundary between the native material and the

weld (fig. 4a). The weld structure was fragmented several times in comparison to the native

material (fig.4b). Solution heat treatment (8h/525oC/air) and ageing (16h/250 ° C / air) of the joint

have resulted in the growth of the grain in the whole joint area. Small and scarce precipitates of

altered morphology were observed in the native material (fig. 4c), the heat affected zone is clear

(fig. 4c) and there is a small phase precipitate in the weld (fig. 4d).

On the basis of the analysis of creep curves of the QE22 alloy welded joints it has been found that

heat treatment improves the creep resistance (fig. 5). The strain of the non heat-treated sample after

110 hours of test was 2.46 % and was over 10 times higher than that of the heat-treated sample after

welding (ε = 0.21%).

The creep rate of the samples tested at the temperature of 180°C increases with the growth of the

stress value (tab. 3). From 5.37*10-9

s-1

to 3.12*10

-8s

-1 for non heat-treated samples and from

2.76*10-11

s-1

to 9.48*10-10

s-1

for heat-treated samples after welding. The creep rates of the non heat-

treated samples, tested at the temperature of 200°C were similar: 4.63* 10-8

s-1

for 60MPa and 4.56*

10-8

s-1

for 70MPa.

Analysis and microstructures at magnifications of up to 200x do not reveal any significant changes

in the weld structure (fig. 6), however, due to the precipitates occurring in the weld after the creep

process, the research should be supplemented with the observations of the substructure. The QE22

alloy joints should be used after the heat treatment.

Based on the performed tests and analysis of their results, the following conclusions have

been drawn:

• The heat treatment of the QE22 alloy joints has an impact on their microstructure. As a

result of solution heat treatment and ageing, the phase precipitates are partly dissolved and

small dispersion precipitates of strengthening phases occur in the microstructure.

• The heat treatment of the QE22 cast magnesium alloy welded joints improves their creep

resistance. The strain of the heat-treated joint after 110 hours of test is several times smaller

than the strain of the non heat-treated joint after welding. The creep rates are also smaller in

case of heat-treated samples.

• The QE22 cast magnesium alloy welded joints repaired with the welding techniques show

creep resistance in the temperature range from 180°C to 200°C. They can work in the stress

conditions up to 90 MPa for the working temperature of 180°C, and up to 70 MPa for the

working temperature of 200°C.

a) b)


Acknowledgements

The study has been financed by the National Science Centre within the project No

2442/B/T02/2011/40 “Structure and properties of welded joints of cast magnesium alloys in

simulated operating conditions”

References

[1] M. Avedesian, H. Baker, Magnesium and Magnesium Alloys. ASM Speciality Handbook,

1999,

[2] H.E. Friedrich, B.L. Mordike : Magnesium Technology: metallurgy, design, data, applications,

Springer-Verlag Berlin Heidelberg, 2006,

[3] A.Szczotok, S. Roskosz : New possibilities of light microscopy research resulting from digital

recording of images, Mat. Sci., 23 (2005) 559-565,

[4] A. Kiełbus: Structure and mechanical properties of casting MSR – B magnesium alloy J. of

Ach. in Mater. and Manuf. Eng., 18 (2006) 131-134,

[5] Elektron MSR-B, Data sheet 463, Magnesium Elektron, Great Britain, 2006,

[6] J. Adamiec , S. Mucha : Determination brittle temperature range of MSR-B magnesium alloy

Arch. of Metal. and Mat, 56 (2011) 117-127,

[7] A. Kierzek , J. Adamiec : Evaluation of susceptibility to hot cracking of magnesium alloy

joints in variable stiffness condition Arch. of Metal. and Mat., 56 (2011), 755-767,

[8] A. Kierzek, J. Adamiec : Design factors influencing weldability of the Mg-4Y-3RE cast

magnesium alloy, 2011 IOP Conf.Ser.: Mater. Sci. Eng, 2011


Microstructure of in situ Mg metal matrix composites based on silica nanoparticles

Anita Olszówka-Myalska1,a, Sam A. McDonald2,b, Philip J. Withers2,c, Hanna Myalska1,d, Grzegorz Moskal1,e

1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice Poland

2 University of Manchester, Grosvenor Street, Manchester M1 7HS, UK

a [email protected], b [email protected], c [email protected], d [email protected], e [email protected]

Keywords: magnesium composite, Mg2Si, nanoparticle silica, differential scanning calorimeter, X-ray tomography

Abstract. Metal matrix composites comprising a magnesium matrix and Mg2Si/MgO dispersoids obtained by hot pressing silica nanoparticle agglomerates and metal powder in a Degussa press were characterized. Two powder mixtures having weight proportions of Mg:SiO2 of 10:0.3 and 10:1 were identically sintered. Their microstructures were characterized by optical microscopy and X-ray diffraction. The size and distribution of the Mg2Si and MgO dispersoids formed in situ were assessed as a function of the original nanosilica content. The behaviour of the composites under compression testing was assessed in 3D by X-ray microtomography using 225kV Nikon X-tek and 150kV Xradia MicroXCT scanners. This provided insights into composite strengthening mechanisms and matrix particle decohesion. Introduction

Magnesium silicide Mg2Si because of its low density (1.91 kg/m3), high melting point (1358 K), good Young’s modulus (120 GPa) and microhardness (100-600 HV), significant coefficient of thermal expansion (7.5x10-6 K-1) [1] is a good candidate as a reinforcing phase in magnesium matrix composites. Further, it can be formed in situ as a reaction product between magnesium and silicon [2-5] or silica [5-8] by a range of technologies. The interphase bonding between Mg2Si and Mg is strong because it forms as a reaction product. However, final material properties depend on the dispersion of Mg2Si in the matrix which is a function of the constituent particulate stock, the consolidation process and subsequent processing. In the case of the silica precursor, various crystalline and amorphous powders feedstocks based on; commercial SiO2 powder with a size range of 3-375µm [5,7], silica rice husks of 3.9-39.2µm [6] and flyash of 100µm [8] have been tried. The aim of this paper is to examine whether agglomerates of nanoparticles of amorphous SiO2 are suitable as a precursor of the Mg2Si phase in magnesium matrix composite obtained by hot pressing, both in terms of initial microstructure and its influence on properties. Methods and results

Material fabrication. The possible reactions between Mg and SiO2 along with their Gibbs free energy evaluated using the HSC Chemistry 4.1 program are presented in Tab.1. These show a negative value of ∆G over the temperature range appropriate for magnesium powder metallurgy processes and confirm the formation of the Mg2Si intermetallic phase as the final reaction product.


Table.1. Thermodynamic data for the possible reactions between Mg and SiO2

Granulated 25-66 µm (99% Aldrich) Mg powder, and nanosized (5nm) silica (aerosil-200, Aldrich) [9] were used which as-manufactured formed the agglomerates (Fig.1). To characterize the reaction scheme during the formation of the Mg2Si differential scanning calorimetry (MULTI HTC SETARAM) was undertaken. The thermal effects during heating and cooling at a rate of 10oC/min in argon atmosphere (99.9999) of a 1:1 powder mixture (by weight) were measured. As seen in fig. 2, during heating a strong exothermic reaction was observed to initiate at 501oC (centred on 515oC) before the endothermic peak associated with the melting point of Mg (centred on 650°C), while Fig.3 shows that during cooling no additional reactions besides crystallisation of Mg were observed.

Fig. 1. SEM micrographs of nanosilica aerosil-200 deposited on graphite conductor and coated with Au conducting film: a) as-manufactured, b) after ultrasonic de-agglomeration.

Fig. 2. DSC trace for the Mg-nanosilica mixture showing exothermic (515°C) and endothermic

(650°C) peaks on heating.

Reaction

Gibbs energy [kJ/mol] 200

oC 500

oC 600

oC 650

oC

4Mg+SiO2→Mg2Si+2MgO -350.78 -337.92 -333.37 -330.99

2Mg+SiO2→2Si+MgO -275.52 -265.38 -261.81 -259.93

2Mg+Si→Mg2Si -75.26 -72.54 -71.57 -71.04


Fig. 3. DSC investigation of Mg-nanosilica mixture during cooling. In order to fabricate sintered composite compacts having a diameter of 20 mm, powder mixtures of magnesium and nanosilica in the weight ratio of 10:0.3 (sample referred to as 2%MgSiO2) and 10:1 (sample referred to as 6%MgSiO2) were prepared. These weight fractions correspond to volume fractions of nanosilica in the powder mixture of approx. 2% and 6%, respectively. Then two hot pressing steps (300oC at 1.5MPa for 10 min and 650oC at 8MPa for 30min) under vacuum in a Degussa press were applied.

Microstructure characterization. Light microscopy (LM) using an Olympus GX71 DP70 was applied to polished composite cross-sections prepared in the presence of water to achieve some contrast. This method is very convenient for Mg2Si phase detection because of the characteristic blue color of magnesium silicide. Microstructure investigations of the 2%MgSiO2 composite showed (Fig. 4) the presence of magnesium grains having an irregular shape typical for pressed metal powder, surrounded by a thin network of dispersed phases containing fine Mg2Si. Fine (1-10µm) Mg2Si inclusions were also occasionally found inside magnesium grains. By contrast, the microstructure of the 6%MgSiO2 composite was quite different (Fig.5). It comprised a multiphase mixture of Mg- Mg2Si/MgO dispersoids having two sizes of Mg2Si grains one 10-50µm (regular) and one 1-3µm (irregular).

Fig. 4. LM micrographs of 2%MgSiO2 composite.

Mg2Si


Fig. 5. LM micrographs of 6%MgSiO2 composite. To characterize the 3D microstructure of the composite compacts, X-ray tomography using a Nikon X-tek 225/320 kV custom bay at an effective pixel size 18.6µm was applied (Fig.6). In both composites pores (black) and particles (white) were detected, and their volume fraction measured (Tab.2). An increase of porosity towards the centre of the coupon was observed whereas the particles are evenly distributed.

Mg2Si


Fig. 6. 3D reconstruction of the a) 2%MgSiO2 and b) 6%MgSiO2 composite compacts showing the distribution of large particles (white) and pores (black) measured on the Nikon X-tek 225/320 KV custom bay.

Table 2. Microstructural parameters extracted from the 3D reconstructions acquired on the, Nikon X-tek X-ray scanner.

Sample

Volume

fractions of

pores

[%]

Volume

fractions of

particles

[%]

Number

of

particles

Largest

particle

volume

[mm3]

Smallest

particle

volume

[mm3]

Mean

particle

volume

[mm3]

2%MgSiO2 0.05 0.021 391 0.012696 0.000019 0.001607

6%MgSiO2 0.27 0.008 151 0.011294 0.000029 0.001197

The phase fractions for the sintered compacts were quantified by X-ray diffraction (XRD) using a JDX-7S diffractometer. According to the JCPDS-International Centre for Diffraction Data 2000 the X-ray diffraction patterns (Fig.7) comprise contributions from Mg, Mg2Si and MgO, however the intensity of Mg2Si and MgO peaks was for the 6%MgSiO2 coupon were higher.


Fig. 7. Profiles of XRD analysis: a) 2%MgSiO2 and b) 6%MgSiO2 compacts.

Characterisation of properties. To measure the density, open porosity and hardness of the sintered composites the Archimedes’ and Brinell indentation (φ=2.5mm, N=250kg) methods were applied. They indicated (Tab. 3) low density and an increase of porosity and hardness with increasing nanosilica content in powder mixture.

Table.3. Properties of the sintered composite compacts.

Cylinders having the dimensions h=8.66 mm, d=3.96 mm for 2%MgSiO2 and h=10.91 mm, d=3.98 mm for 6%MgSiO2 were cut in order to characterise the compressive behaviour. These cylinders were imaged using an Xradia MicroXCT system with effective pixel size 5.64µm collecting 1299-1948 projections. The Avizo Standard programme was applied to segment the phases in the resulting 3D reconstructed volumes. The same procedure was repeated after the compression tests which were carried out using a MTS Alliance RT/100 test machine (Fig.8.). The 3D images of samples before and after compression test are presented in figures 9-12. They confirm the presence of micropores and particles registered earlier in both materials but most importantly they demonstrate the different manner in which crack defects propagate in the two composites. For 2%MgSiO2 a lot of small cracks were formed and some degree of micropore closing was evident while for 6%MgSiO2 only a few (four) large cracks formed, mainly having a shear orientation of 45o to the compression axis. Moreover, in the sample 6%MgSiO2 an increase of the initial pores size as a result of compression test was registered.

Material Density

[g/cm3] Porosity [%] (Archimedes)

Porosity [%] (X-ray tomography)

Hardness HB

2%MgSiO2 1.75 0.09 0.05 47.5±0.1

6%MgSiO2 1.82 0.32 0.27 54.3±8.7


Fig. 8. Profile of force vs. deformation for the compression tests corresponding to the images below.

Fig. 9. 3D Visualization of the 2%MgSiO2 test-piece: a) as-manufactured, b) after compression testing recorded on the Xradia MicroXCT.

Plane XY Plane XZ

Fig. 10. 3D Visualization of cracks (colour) formed in 2%MgSiO2 composite after compression testing, shown superimposed on a tomographic slice (black and white) Xradia MicroXCT.


Fig. 11. 3D Visualization of the 6%MgSiO2 test-piece: a) as-manufactured, b) after compression testing recorded on the Xradia MicroXCT.

Plane XY Plane XZ

Fig. 12. 3D Visualization of cracks (colour) in 6%MgSiO2 composite after compression testing, shown superimposed on a tomographic slice (black and white) Xradia MicroXCT.

Discussion

Our observations confirm the reaction between nanosilica and magnesium during sintering, revealed by the strong exothermic peak in the DSC and by x-ray diffraction peaks corresponding to Mg2Si and MgO. The reaction peaks at 514oC in agreement with Umeda et al. [6] for amorphous silica (3.9 and 6.8µm granulates). As one might expect, the XRD profiles reveal an increase in Mg2Si and MgO with increasing nanosilica content in the initial powder mixture. Our observations of Mg2Si formation for nanoparticles are in good agreement with literature data for micro-sized silica particles [5-8].

Significant differences were observed according to the original nanosilica content. For the 2%MgSiO2 compact, globular Mg grains were surrounded by a thin, multiphase network containing very fine irregular grains of Mg2Si. The 6%MgSiO2 material was a relatively homogenous mixture of phases containing micro-sized Mg2Si grains. For both samples some areas showed the presence of Mg2Si nanosized grains mixed with MgO. While optical microscopy is very convenient method for Mg2Si phase detection (blue color of silicide), this is not sufficient at such high resolutions. Similarly, unpublished results suggest SEM observations using SEI and BSE techniques are not helpful either. Therefore in further work electron transmission microscopy (TEM) with selected area diffraction patterns (SADP) will be necessary.


Results of X-ray tomography of the as-manufactured composite compacts showed a characteristic uniform distribution of the micropores typical for compacted bars. Generally, the porosity of composite was low, being a little bit higher for sample 6%MgSiO2 (0.27%) than for 2%MgSiO2 (0.05%). This tendency was in good agreement with porosity measurements by Archimedes’ method. In addition to the pores the X-ray tomography detected the presence of 390 and 150 light element containing particles having a diameter of 200-300µm in the 2%MgSiO2 and 6%MgSiO2 samples respectively. This observation needs to be interpreted in the light of the LM observations. This would suggest that agglomerates of the Mg2Si/MgO dipersoids are being picked up in X-ray images. A strengthening effect as a result of the increasing Mg2Si/MgO dispersoids volume fraction was confirmed by the hardness measurements and was evident in the associated compression curves. X-ray microtomography revealed the manner in which different starting microstructures lead to different damage accumulation processes. For the 2%MgSiO2 sample the pore fraction was smaller leading to the initiation of a lot of small cracks pores and reasonable plasticity. By contrast, the 6%MgSiO2 composite the higher pore fraction lead to pore growth, reduced plasticity and the growth of a few catastrophic cracks at lower strains. Conclusions • Nanosilica agglomerates are a suitable starting constituent for the fabrication of magnesium matrix composite forming Mg2Si/MgO dispersoids forming a network at the boundary of the magnesium grains or a homogenous multiphase mixture. • The Mg2Si was less than 50µm for the 6%MgSiO2 sample and less than 10µm for the 2%MgSiO2 sample. TEM studies are needed to verify the probable presence of nanosized Mg2Si particles. • With increasing dispersoid fraction the hardness and compressive strength increased. • X-ray tomography examination of the Mg-Mg2Si/MgO composite indicated that the porosity was greater for the high particle fraction composite; in addition agglomerates were observed that were not identified by LM observations. • The X-ray microtomograhy studies revealed differences in composite cracking under compressive straining that differed according to the starting microstructure.

Acknowledgement

Funding is acknowledged through EPSRC grants EP/F007906, EP/F028431 and EP/I02249X to set up and maintain the Henry Moseley X-ray imaging Facility.

References

[1] C.C. Koch Nanostructur. Mater., The synthesis and structure of nanocrystalline materials produced by mechanical attrition: A review, Nanostructur. Mater. 2 (1993) 109-129. [2] L.Lu, K.K. Thong., M. Gupta, Mg-based composite reinforced by Mg2Si, Comp. Sci. and Techn. 63 (2003) 627-632. [3] K. Chen, Z.Q. Li, J.S. Liu, J.N. Yang, Y.D. Sun, S.G. Bin, The effect of Ba addition on microstructure of in situ synthesized Mg2Si/Mg-Zn-Si composites, J. of Alloys and Comp. 487 (2009) 293-297. [4] Z. Trojanová, V. Gärtnerová, A. Jäger, A. Námešný, M. Chalupová, P. Plaček, P. Lukáč, Mechanical and fracture properties of an AZ91 alloy reinforced by Si and SiC particles, Comp. Sci. and Techn. 69 (2009) 2256-2264. [5] K. Kondoh, H. Oginuma and T. Aizawa, Tribological Properties of Magnesium Composite Alloy with In situ Synthesized Mg2Si Dispersoids, Mater Trans., 44 (2003) 524-530.


[6] J. Umeda, K. Kondoh, M. Kawakami, H. Imai, Powder metallurgy magnesium composite with magnesium silicide in using rice husk silica particles, Powder Techn. 189 (2009) 399-403. [7] M. Aydin, C. Özgür, O.San, Microstructure and hardness of Mg-based composites reinforced with Mg2Si particles, Rare Met. 28 (2009) 396-400. [8] Z. Huang, S. Yu: Microstructure characterization on the formation of in situ Mg2Si and MgO reinforcements in AZ91D/Flyash composites, J. of Alloys and Comp. 509 (2011) 311-315. [9] M. Sopicka-Lizer, R.A. Terpstra, R. Metselaar: Carbothermal production of beta'-sialon from alumina, silica and carbon mixture", J. Mater. Sci. 30, (1995) 6363-6369.


Microstructure of Mg-Ti-Al composite hot pressed at different temperature

Anita Olszówka-Myalska 1,a, Roman Przeliorz 1,b, Tomasz Rzychoń 1,c,

Monika Misiowiec 1,d


a [email protected], b [email protected], c tomasz.rzychoń@polsl.pl, d [email protected]

Keywords: magnesium composite, titanium particles, Mg17Al12, Al3Ti, differential scanning calorimeter

Abstract. Metal matrix composite comprising a multiphase magnesium matrix and titanium

particles fabricated by hot pressing was characterized. Powder mixture of the Mg:Ti:Al at weight

ratio equal to 10.5:6.1:3.4 was sintered at 640, 650 and 660

oC whereas other parameters were held

constant. Thermal effects during heating and cooling of powder mixture were measured by

differential scanning calorimetry (DSC). Microstructure of composite was characterized by

scanning electron microscopy (SEM) with a use of X-ray energy dispersive spectroscopy (EDS) and

X-ray diffraction (XRD). For all conditions of components consolidation α-Mg, α-Ti, Mg17Al12 and

Al3Ti were identified. It was revealed that dispersion and location of Mg17Al12 and Al3Ti

compounds depended on sintering temperature. Measurements of hardness and density of obtained

non-porous composite gave approximate results of 130 HV and 2.7 g/cm3 respectively.

Introduction An application of titanium powder in magnesium matrix composite is discussed in some recently

published works because of metallic components low density and good mechanical properties of

titanium. Generally, two ideas of titanium particles addition into magnesium matrix are discussed.

Titanium can be a reinforcement in Mg-(Ti)p ex situ composite [1,2] because of titanium inactivity

and good wettability with magnesium or can be a precursor of A3Ti titanium aluminide in in situ

composite. In the case of the titanium aluminide formation in magnesium matrix, the technologies

employing aluminium addition as separated particles [3] or as an alloying element in magnesium

based alloy [4] were reported.

Generally, the titanium aluminide in situ formation is described by two different mechanisms, a

relatively slow growth controlled by the diffusion (when dispersion of a new phase depends on

titanium powder dispersion) or by combustion synthesis (when very fine particles of a new phase

are obtained [5]). However, the combustion synthesis of titanium aluminides is strongly exothermal

and may result in the equipment damage and/or to generate a porous microstructure of a composite

fabricated by powder metallurgy processes (PM). Different PM technological procedures like cold

pressing, sintering and additional annealing [3], cold pressing, sintering and extruding [4] are

proposed in literature for magnesium matrix composite with titanium aluminides.

In the experiment presented in this paper a pressure sintering of conventionally mixed Mg, Ti and

Al powders (in nitrogen atmosphere) was proposed as a relatively simple and effective method. The

aim of the paper was to show an effect of hot pressing temperature (from the range of 640-660oC)

on final microstructure and some properties of composite obtained from powder mixture of Mg, Ti

and Al at the weight ratio of 10.5:6.1:3.4.


Experimental procedure

Materials and methods

The starting materials used to prepare the powder mixtures by milling in a Fritsch mill in nitrogen

atmosphere were Mg (particles 25-66 µm, 99%, Sigma Aldrich), Ti (particles, 44 µm, 99,9%

,Sigma Aldrich) and Al (flakes, 25-100 µm, 99.7%, Benda-Lutz Skawina). Their 3D micrographs

obtained with scanning electron microscope (SEM) are presented in Fig. 1.

In order to determine the thermal effects occurring during heating and cooling of powder mixture

Mg-Al-Ti (at the weight ratio of 10.5:6.1:3.4) and for comparison of Mg-Ti mixture (at the weight

ratio of 10:6) the differential scanning calorimetry (MULTI HTC SETARAM) was employed. The

experiment was carried out in argon atmosphere (99.999) at the heating/cooling rate of 10oC/min

for temperature range of 20-700oC.

The Degussa press was applied for the consolidation process of Mg-Ti-Al mixture (weight ratio of

10.5:6.1:3.4) under vacuum of 2.8Pa in two hot pressing steps. At first 30 minutes under the

pressure of 15MPa at the temperature of 300oC and then 30 minutes under the pressure of 8MPa at

the temperature of 640oC, 650

oC and 660

oC. Respective samples are referred as MgTiAl-1,

MgTiAl-2 and MgTiAl-3.

To measure the density and open porosity the Archimedes’ method was applied whereas Vickers

indentation (Duramin A, HV3, HV0.1) was employed for hardness determination of the sintered

composites.

The microstructure and composition of polished and unetched cross sections was characterized by

scanning electron microscopy (SEM Hitachi 4200) with energy dispersive spectroscopy (EDS,

Noran system). Quantitative description of composite microstructure focusing mainly at titanium

particles features was obtained with the Metilo software [6].

Phase constitution of manufactured sinters was examined by X-ray diffraction method (XRD, JEOL

JDX-7S diffractometer, JCPDS-International Centre for Diffraction Data 2000).

Fig. 1. SEM micrographs of metal powders applied in experiment: a) Mg, b) Ti, c) Al



Thermal analysis

The DSC profile of Mg-Ti mixture (Fig.2a) exhibited only one strong and sharp endothermic effect

with onset point at 646.6oC and peak max. at 658.2

oC at heating and two exothermic effects at

cooling (one strong and sharp with onset point at 645.4oC and peak max. at 634.1

oC, and another

weak at 447oC). The DSC plot suggests only melting and crystallization of magnesium and

confirms an absence of Mg-Ti chemical interaction at the experiment temperature.

In the case of Mg-Ti-Al system the DSC profile (Fig.2b) exhibits some differences in comparison

to Mg-Ti system. One strong endothermic effect (with onset point at 515.6oC and peak max. at

541.5oC) followed by an exothermic effect (with onset point at 607.4

oC and peak max. at 638.2

oC)

were determined at heating. Two exothermic effects (one strong with onset point at 629.5oC and

peak max. at 612oC and one weak with onset point at 433.7

oC and peak max. at 423.4

oC) at cooling

were observed. Combination of the new peak appearance (exothermic effect) and the endothermic

peak broadening/reduction/shift seems to be the result of Al-Ti exothermic reaction. Comparison of

DSC profiles obtained at cooling shows some differences of a strong exothermic peak (broadening,

intensity reduction and shift). It can be explained by a decrease of liquid metal in system as a result

of magnesium matrix enrichment with other elements and phases.

Fig. 2. Profiles of DSC obtained for powder mixtures of : a) Mg+Ti, b) Mg+Ti+Al


Microstructure

X-ray diffraction analysis.

The X-ray diffraction patterns (Fig.3) show the presence of α-Mg, α-Ti, Mg17Al12 and Al3Ti phases

in investigated material, independently of the applied sintering temperature. Aluminium residuals

were not detected. That result suggests formation of composite with the α-Mg/Mg17Al12 multiphase

matrix typical for AZ magnesium alloys and α-Ti particles. The Al3Ti intermetallic compound

location will be discussed in the next section (EDS). In the light of thermodynamic data (Gibbs free

energy value varying from -12 kJ/mol to -30 kJ/mol at the temperature range of 300- 650oC [7]) the

presence of Mg17Al12 compound is reasonable however not confirmed by authors of work [3].

Although Al3Ti formation is more probable (∆G value approx. -34 and -31 kJ/mol at the

temperature range of 300-650oC respectively [8]) but the powder mixture consolidation conditions

ensure a direct contact of Mg-Al and Ti-Al particles and following formation of both intermetallic

phases by diffusion. For quantitative characteristics of phase composition changes with sintering

temperature further investigations are necessary.

Fig. 3. XRD patterns of Mg-Ti-Al composite

Microscopic investigations.

Results of SEM observations of composites cross sections combined with X-ray mapping of Mg, Ti

and Al are presented in Figures 4-9. It is shown that in all samples titanium particles (bright) are

surrounded with the matrix containing α-magnesium (dark) and irregular areas (grey). X-ray

mappings for composite sample MgTiAl-1 (Fig.4. ) show enrichment of aluminium in grey areas

(likely Mg17Al12) in comparison to their content in a dark matrix (α-Mg). In case of next samples

MgTiAl-2 (Fig.6.) and MgTiAl-3 (Fig.8.) Al redistribution from grey areas to matrix/Ti-particle

interface is observed. Higher magnification mappings (Fig.5,7,9) confirm Al redistribution and

evident increase of its content around titanium particles with some amount of titanium in the same

microareas. That suggests location of Al3Ti phase in a continuous zone with thickness less than

0.5µm (sample MgTiAl-1) on Ti particles and fine Al3Ti particles in the vicinity of Ti particles

(MgTiAl-2 and MgTiAl-3).

Quantitative characteristics of Ti particles by image analysis (Tab.1) show a reduction of their mean

area, perimeter, diameter and shape coefficient being a result of titanium reaction in Mg-Ti-Al

system. Moreover, the micro-hardness measurements (Tab.2) reveal an increase of Ti particles


hardness with a sintering temperature increase as an effect of α-Ti solid solution formation.

Explanation of similar changes of matrix micro-hardness can be expressed in terms of diffusion

processes followed by microstructure modification.

Obtained results are in good agreement with mechanism of Al3Ti formation presented in [3] for

liquid Mg-Al alloys-solid Ti system, however a combustion synthesis with full titanium

consumption did not appear at applied conditions. A presence of Mg17Al12 phase in the samples

sintered at the temperature close to magnesium melting point indicates aluminium partition between

Mg-Al and Ti-Al reactions. Only significant increase of hot pressing time at 650-660oC can be

supposed to intensify consumption of Al from Mg-Al alloy.

Fig.4. SEM micrograph of MgTiAl-1 composite and X-ray mapping of Mg, Ti and Al


Fig.5. SEM micrograph of single Ti particle in MgTiAl-1 composite and X-ray mapping of Mg,

Ti and Al

Fig.6. SEM micrograph of MgTiAl-2 composite and X-ray mapping of Mg, Ti and Al.



Ti and Al

Fig.8. SEM micrograph of MgTiAl-3 composite and X-ray mapping of Mg, Ti and Al



Ti and Al

Table 1. Quantitative characteristics of Ti particles in composite.

Sample MgTiAl-1 MgTiAl-2 MgTiAl-3

Number of characterised

particles 51 53 58

Mean perimeter, µm 1018,37 889,83 623,97

Shape coefficient 0,48 0,43 0,43

Maximal diameter, µm 244,10 208,20 153,07

Minimal diameter, µm 149,93 125,37 93,73

Properties characterization

The measurements of density and open porosity of fabricated composite samples indicate (Tab.2)

low porosity (Archimedes’ zero) and density similar to that of aluminium alloys. Hardness of

composite samples was very high (130HV) comparing to commercial magnesium alloys (approx.

70HV). Additionally, Ti particles (strengthened solid-solution) and matrix microhardness is

growing with a sintering temperature. This is the confirmation of diffusion effect on final

microstructure examined by SEM+EDS methods. In order to explain insignificant hardness

decrease of MgTiAl-3 sample in comparison to AlTiAl-2 sample hardness the nanoindentation

seems to be necessary.


Table.2. Results of Archimedes, and hardness measurements.

Sample

Sintering

temp.

[oC]

Density

[g/cm3]

Open

porosity

[%]

Hardness

HV3

Microhardness HV0.1

Ti particle matrix

MgTiAl-1 640 2.63 0 128.8±9.3 178.5±23 91±25

MgTiAl-2 650 2.68 0 137.1±6 196±32 124±22

MgTiAl-3 660 2.66 0 130.1±5.8 215.5±18 149±1

Conclusions

• Composite material with Ti particles and multiphase magnesium base matrix of hardness 130HV

and density 2.7g/cm3 was obtained by hot pressing of a conventionally mixed Mg-Ti-Al powders at

the temperature range of 640-650oC.

• Microstructure investigations showed complete consumption of aluminium and decrease of

titanium particles initial size.

• Two types of aluminides Al3Ti and Mg17Al12, and solid solutions of α-Ti and α-Mg were

identified by XRD method in composite samples. The Mg17Al12 phase was formed as a dispersoids

in magnesium matrix and it dispersion increased with sintering temperature. The Al3Ti aluminide

was formed either as a thin layer around Ti particles for process at 640oC or as a dispersoids in

magnesium matrix at higher temperature but with a tendency of concentration around Ti particles.

• An application of Mg-Al powder mixture for in situ synthesis of Al3Ti phase by powder

metallurgy processing at the temperature close to magnesium melting point results in formation of

Mg17Al12 phase in magnesium matrix.

References

[1] L. Lu, M.O., Lai, L. Froyen, Effects of mechanical milling on the properties of Mg-10.3%Ti

and Mg-5%Al-10.3%Ti metal-metal composite, J. of Alloys and Comp. 382, (2005) 260-264.

[2] J. Umeda, M. Kawakami, K. Kondoh, H. Imai, Microstructural and mechanical properties of

titanium particulate reinforced magnesium composite materials, Mater. Chem. and Phys., 130

(2010) 649-657.

[3] Yang Z.R, Wang S.Q., Cui X. H., Zhao Y.T., GaoM.j, Wei M.X., Formation of Al3Ti/Mg

composite by powder metallurgy of Mg-Al-Ti system, IOP Sci. Technol. Adv. Mater. 9 (2008)

1-6.

[4] B.K. Raghunath, R. Karthikeyan, G. Ganesan, M. Gupta, An investigation of hot deformation

response of particulate-reinforced magnesium + 9% titanium composite, Mater. and Design, 29

(2008) 622-627.

[5] A. Olszówka-Myalska, Evolution of titanium particles microstructure in aluminium matrix

composite obtained by powder metallurgy method, Inżynieria Materiałowa, 3-4 (2007) 200-

203.

[6] J. Szala, Application of computer-aided image analysis methods for a quantitative evaluation of

material structure (in Polish), Silesian Technical University, 2001 Gliwice-Poland.

[7] H. Zhang, S.L. Shang, Y. Wang, A. Saengdeejing, L.Q. Chen, Z.K. Liu, First-principles

calulations of the elastic, phonon and the thermodynamic properties of Al12Mg17, Acta Mat. 58

(2010) 4012-4018

[8] M. Sujata, S. Bhargava, S. Sangal, On the formation of TiAl3 between solid Ti and liquid Al, J.

of Mat. Sci. Lett. 16 (1997) 1175-1178


CHAPTER 3:

Titanium and Titanium Alloys

The chemical composition and microstructure of Ti-47Al-2W-0.5Si alloy melted in ceramic crucibles

Wojciech Szkliniarz 1,a, Agnieszka Szkliniarz 1,b



Keywords: TiAl-based alloys, melting, ceramic crucibles.

Abstract. In this paper, the technology of melting in induction furnaces with ceramic crucibles was used for production of TiAl-based Ti-47Al-2W-0.5Si alloy. Due to high reactivity of liquid titanium alloys, the melting process was conducted in special crucibles made of stabilised ceramic materials resistant to the aggressive action of these alloys. When characterising the chemical composition and microstructure of Ti-47Al-2W-0.5Si alloy melted in different ceramic crucibles, problems accompanying the melting process were described and conditions for making an alloy with satisfactory purity were determined.

Introduction

TiAl intermetallic alloys with high melting point, low density and good oxidation and creep resistance can be used for components of aircraft and car engines as well as frame and sheathing of space shuttles elements to operate within the range of temperatures from 600 to 850°C [1÷5]. The highest expectations concern the use of such alloys as substitutes for expensive and heavy nickel superalloys for manufacturing low-pressure turbine blades and high-pressure compressor rotor blades [2÷4].

TiAl-based alloys are most often produced by melting and casting in the arc, electron-beam, plasma and induction furnaces in vacuum or argon atmosphere [3]. The cold-wall induction melting in water-cooled copper crucibles is used for melting them. Local charge melting and lack of efficient stirring within the entire alloy volume during the arc, electron-beam and plasma melting result in large variations in chemical composition and microstructure of produced ingots or castings. The technology of induction melting in water-cooled copper crucibles is devoid of the above-mentioned faults, although the intensive crucible cooling makes it difficult to obtain the required alloy overheating, necessary for correct course of the casting process [6].

For this reason, more and more often the trials to use the technology of melting in induction furnaces with ceramic crucibles, which is widely used in manufacturing alloys of other metals, for production of TiAl-based alloys are made. The drawback to use this technology for melting TiAl-based alloys is their high reactivity with the ceramic crucible, which may result in alloys contamination as a result of chemical reactions between liquid alloy and ceramic materials of the crucible. High process temperature, exothermal reactions accompanying the melting process and intensive electromagnetic stirring additionally promote the degradation process of the crucible and uncontrolled passing of its components into liquid alloy Generally, the chemical composition, purity and microstructure of TiAl-based alloys melted in ceramic crucibles in vacuum induction furnaces depend on the crucible material, furnace atmosphere, melting temperature and time well as the form and purity of charge materials used in the process [3, 7, 8].

The investigations carried out so far indicated that the most suitable ceramic crucibles for melting TiAl-based alloys were made of CaO [9÷13]. The melting in such crucibles, much cheaper as compared to melting in water-cooled copper crucibles, was first used in Japan and Korea for production of lightweight car valves made of TiAl-based alloys [12, 13]. There has also appeared the information on possibility of using crucibles made of CaO with 5% additive of CaF2 [14] and ceramic crucibles made of conventional materials with protective layer applied to their internal


surface, resistant to liquid titanium alloys [15], for melting TiAl-based alloys. The latter ones are much cheaper than solid crucibles made of CaO. They also not subject to fast degradation as a result of hydration effected by humidity in the air, and thus less troublesome when stored.

The main aim of the investigations was assessment of the possibility to melt Ti-47Al-2W-0.5Si alloy in vacuum induction furnaces with ceramic crucibles. When characterising the degree of degradation of individual crucibles after melting as well as the chemical composition and microstructure of alloy melted in them, problems accompanying the melting process were described and conditions for making an alloy with satisfactory purity were determined.

Materials and Research Methodology

The material for investigations was Ti-47Al-2W-0.5Si alloy with nominal chemical composition as specified in Table 1, belonging to the group of TiAl intermetallic alloys.

Table 1. Nominal alloy composition

Element Al W Si Ti [at.%] 47.00 2.00 0.50

Balance [wt.%] 31.18 9.04 0.35

Alloy was melted in Balzers VSG-02 and Leybold-Heraeus IS-5/III vacuum induction furnaces,

in special ceramic crucibles made of different materials. The crucibles made of oxide ceramic materials and characterised by lower standard free formation energy than that of TiO2 were mainly used [9, 10]. The capacity of the crucibles ranged from approx. 0.2 to 3.5 l. The alloy melting process included: preparation of charge materials and casting mould, preparation of crucible and it installation in the induction furnace coil, melting and casting into the cold graphite mould. The charge in the form of technically pure titanium and Al-W-Si master alloy was used. The melting process was conducted in the atmosphere of highest purity argon with pressure of approx. 0.08 MPa at 1650°C. The obtained ingots of 30 mm in diameter were subject to homogenising at 1400°C for 1 h. Then the ingots were cooled with the furnace.

In the obtained Ti-47Al-2W-0.5Si alloy ingots, the content of basic alloy components (aluminium, tungsten, silicon), oxygen and other impurities produced as a result of reaction between liquid alloy and crucible material as well as macro- and microstructure were analysed. The content of oxygen was determined by high-temperature extraction method and of carbon – by HFIR method. The analysis of remaining components content was made based on the OES-ICP method. The samples for macro- and microstructure investigations were etched with Kroll reagent. The microstructure investigations were carried out with Nicon Epiphot 200 optic microscope and Hitachi S-4200 scanning microscope equipped with EDS detector Voyager of Noran Instruments.

Research Results

The possibility of melting a TiAl-based alloy with satisfactory purity in the vacuum induction furnaces with ceramic crucibles should be sought in the minimisation of the aggressive impact of liquid alloy on crucible materials. It was assumed that this could be obtained by the selection of the appropriate crucible material, use of argon rather than vacuum shield and use of master alloy rather than pure components as the charge. It should result in reduction in intensity of chemical reactions between the crucible material and liquid alloy and minimisation of melting temperature and duration. This, in turn, should contribute to the reduction in the mass of substances passing from the melting space into liquid alloy, which results in its contamination with elements contained in the crucible material.


Every ingot melted in the induction furnace with ceramic crucible, under argon shield and with use of the master alloy as a charge was characterised by good surface quality (Fig. 1a). In the coarse-grained macrostructure (Fig. 1b) of the ingot made from Ti-47Al-2W-0.5Si alloy after homogenisation, there is a wide zone of columnar crystals in the external part of the ingot and much smaller zone of equiaxed crystals inside the ingot.

a) b) c)

Fig. 1. Ingot (a), its macrostructure (b) and expected microstructure (c)

The typical expected alloy microstructure after casting and homogenising should be characterised by fine-grained structure with alternately arranged lamellar precipitations of α2 and γ phases inside the grains (Fig. 1c). The typical, although coarse-grained, lamellar microstructure is observed in alloy melted in CaO crucible (Fig. 2a), while lamellar microstructures of alloys melted in other crucibles contain phase precipitations due to the reactive nature of liquid alloy in relation to the ceramic crucible material (Fig. 2b-d).

a) b)

c) d)

Fig. 2. Microstructure of alloy melted in CaO (a), graphite (b), SiC (c) and ZrO2 (d) crucible after homogenisation


The analysis of chemical composition of Ti-47Al-2W-0.5Si alloy melted in different crucibles (Tab. 2) shows that in addition to the basic components (aluminium, tungsten, silicon) there is also oxygen and other components that were originally included in the material of particular crucibles.

Table 2. Chemical composition of Ti-47Al-2W-0.5Si alloy melted in ceramic crucibles

Ceramic crucibles Alloy composition, [wt.%]

Al W Si O Others ZrO2 30.61 11.05 0.33 0.79 Zr: 3.72 MgO 31.34 10.54 0.28 0.58 Mg: 0.12 CaO 30.77 10.07 0.21 0.06 Ca: 0.09 SiC 32.47 10.41 0.45 0.21 C: 0.21 Graphite Lack of possibility to cast alloy Isostatic graphite 31.34 11.51 0.29 0.08 C: 0.80 High-density isostatic graphite 31.75 10.00 0.35 0.06 C: 0.13

ZrO2 crucible

Alloy melting in ZrO2 crucible is accompanied by chemical reactions at the liquid alloy/crucible interface the visible effect of which is steaming, ejection of liquid alloy from crucible and remains of solidified alloy on the internal crucible surface after melting (Fig. 3a). The degradation of the internal crucible surface results in passing the basic crucible component, i.e. ZrO2, to liquid alloy where it is decomposed as a result of reaction:

ZrO2 (crucible) = Zr (liquid alloy) + 2O (liquid alloy) (1)

a) b)

c) d)

Fig. 3. ZrO2 crucible after first melting (a), microstructure of alloy melted in it (b) and chemical composition of precipitation visible in Fig. 3b (c) and of lamellar matrix (d)


The effect of these processes is very high content of zirconium and oxygen (Tab. 2), which are components of the crucible material, in alloy. The existence of large ZrO2 particles (Fig. 3b, c) in two-phase lamellar microstructure of alloy and zirconium in metallic matrix (Fig. 3b, c) shows that the reaction (1) is not complete.

MgO crucible

The visible indication of chemical reactions occurring at the liquid alloy/crucible interface during alloy melting in MgO crucible is the effects of strong steaming and remains of solidified alloy on the crucible’s bottom and internal wall from the pouring side, ended with a characteristic “tongue” in the upper part of the crucible (Fig. 4a). Alloy melted in MgO crucible shows the typical lamellar microstructure (Fig. 4b, c) after casting and homogenising, and its chemical composition reveals high oxygen content and the existence of magnesium (Tab. 2) as a result of passing the basic crucible component, i.e. MgO, into liquid alloy and its decomposition as a result of reaction:

MgO (crucible) = Mg (liquid alloy) + O (liquid alloy) (2)

The presence of single MgO particles in alloy microstructure (Fig. 4c, d) shows that also in this case the reaction (2) was not complete.

a) b)

c) d)

Fig. 4. MgO crucible after first melting (a), microstructure of alloy melted in it (b, c) and chemical composition of precipitation visible in Fig. 4c (d)

CaO crucible

The effects of strong steaming accompany Ti-47Al-2W-0.5Si alloy melting in CaO crucible. It does not influence on very good condition of crucible after melting where only a slight remain of solidified alloy was found on the crucible’s bottom (Fig. 5a). The condition of CaO crucible after


melting and the presence of exceptionally low oxygen content and low calcium content (Tab. 2) in chemical composition of alloy melted in it prove that passing the basic crucible component, i.e. CaO, into liquid alloy and its decomposition as a result of reaction:

CaO (crucible) = Ca (liquid alloy) + O (liquid alloy) (3)

take place with low intensity. Alloy melted in CaO crucible has lamellar microstructure with strongly fragmented α2 phase

(Fig. 5b, c) after casting and homogenising. Calcium, the existence of which was found in chemical composition of this alloy, occurs in the form of single particles with regular shapes (Fig. 5c, d).

a) b)

c) d)

Fig. 5. CaO crucible after first melting (a), microstructure of alloy melted in it (b, c) and chemical composition of precipitation visible in Fig. 5c (d)

SiC crucible

No adverse effects accompany alloy melting in SiC crucibles. However, the appearance of crucible with remains of solidified alloy on its bottom and internal side walls after melting (Fig. 6a) as well as increased carbon and silicon content in alloy (Tab. 2) show that also in this case there are intensive reactions at the liquid alloy/crucible interface. The result of crucible degradation, passing of the basic crucible component, i.e. SiC, into liquid alloy, its decomposition and reactions between products of this decomposition and liquid alloy:

SiC (crucible) = Si (liquid alloy) + C (liquid alloy) (4a)

3Si (liquid alloy) + 5Ti (liquid alloy) = Ti5Si3 (liquid alloy) (4b)


is increased carbon and silicon content in alloy (Tab. 2) and the occurrence of numerous Ti5Si3 phase precipitations (Fig. 6b, c) in its microstructure, whereas in the metallic matrix only the basic alloy components were found (Fig. 6d).

a) b)

c) d)

Fig. 6. SiC crucible after first melting (a), microstructure of alloy melted in it (b) and chemical composition of precipitations visible in Fig. 6b (c) and of metallic matrix (d)

Graphite crucible

The attempt to melt alloy in a graphite crucible was a complete failure due to the lack of possibility to pour it from crucible (Fig. 7). The reason is very high alloy viscosity caused by uncontrolled passing of the basic crucible component, i.e. carbon, into liquid alloy and high affinity between carbon and titanium.

Fig. 7. Graphite crucible after first melting


Isostatic graphite crucible

Although it does not present any problems with pouring of liquid alloy from crucible (Fig. 8a), Ti-47Al-2W-0.5Si alloy melting in isostatic graphite crucible does not allow obtaining an alloy with satisfactory purity the proof of which is high carbon content in the obtained alloy (Tab. 2). The remains of solidified alloy on the crucible’s bottom and internal wall from the pouring side, ended with a characteristic “tongue” in the upper part of the crucible (Fig. 8a), indicate that also in this case there are intensive reactions at the liquid alloy/crucible interface. After casting and homogenising, alloy shows the characteristic lamellar microstructure with fragmented α2 phase precipitations (Fig. 8b). Carbon passing from crucible into liquid alloy results in existence of fine and large precipitations (Fig. 8b, c) with increased carbon content (Fig. 8d) in its lamellar microstructure.

a) b)

c) d)

Fig. 8. Isostatic graphite crucible after first melting (a), microstructure of alloy melted in it (b, c) and chemical composition of large precipitations visible in Fig. 8c (d)

High-density isostatic graphite crucible

Ti-47Al-2W-0.5Si alloy melted in high-density and purity isostatic graphite crucible is characterised by very low oxygen content and carbon content acceptable for this group of alloys (Tab. 2). The condition of the internal crucible surface after melting (Fig. 9a) shows there are no reactions at the liquid alloy/crucible interface. Also in the lamellar microstructure of alloy melted in this crucible no existence of carbon-rich precipitations was found (Fig. 9b).


a) b)

Fig. 9. High-density isostatic graphite crucible after first melting (a) and microstructure of alloy melted in it (b)

Summary

The investigations have revealed the possibility of melting Ti-47Al-2W-0.5Si alloy with assumed chemical composition and satisfactory purity in vacuum induction furnaces with CaO or high-density isostatic graphite crucibles. The CaO crucibles are very unstable. Due to high hygroscopicity, they require special storage conditions. Melting in these crucibles is accompanied by strong steaming.

High reactivity of liquid alloy causes crucible degradation as a result of chemical reactions between liquid alloy and ceramic materials of crucibles. They are accompanied by strong steaming and ejection of liquid alloy from crucible during melting as well as the remains of solidified alloy on the crucible’s bottom and internal side walls, sometimes ended with a characteristic “tongue” in the upper part of the crucible, from the pouring side.

The crucible degradation is accompanied by uncontrolled passing of components originally included in crucible material into melted alloy and their partial or complete decomposition, and even their reactions with liquid alloy components. It usually results in unacceptable content of oxygen and other crucible components in alloy melted in ceramic oxide crucibles, silicon and carbon in alloy melted in SiC crucibles, and carbon in alloy melted in graphite crucibles.

The products of partial and complete decomposition of crucibles components are in the cast alloy in an undecomposed form, in the form of secondary solid solutions based on intermetallic phases occurring in alloy, in a free form and in the form of other intermetallic phases formed with liquid alloy components. Their presence usually has a modification effect on alloy microstructure, which consists of alternately arranged lamellar precipitations of γ and α2 phases, thus resulting in fragmentation of lamellar precipitations of α2 phase.

Acknowledgment

This scientific work is financed from the budget funds for science in the years 2010-2013 as the research project no NR15-0019-10


References

[1] E. A. Loria, Gamma titanium aluminides as prospective structural materials, Intermetallics 8 (2000) 1339-1345.

[2] X. Wu, Review of alloy and process development of TiAl alloys, Intermetallics 14 (2006) 1114-1122.

[3] W. Szkliniarz, Stopy na osnowie faz międzymetalicznych z układu Ti-Al, Wyd. Pol. Śl., Gliwice, 2007.

[4] A. Lasalmonie, Intermetallics: Why is it so difficult to introduce them in gas turbine engines?, Intermetallics 14 (2006) 1123-1129.

[5] H. Clemens, H. Kestler, Processing and Applications of Intermetallic γ−TiAl-Based Alloys, Advanced Engineering Materials 9 (2000) 551-570.

[6] G. Jarczyk, M. Blum, P. Busse, H. Scholz, H.-J. Laudenberg, K. Segtrop, New casting technology for low-priced titanium-aluminide automotive valves, Inżynieria Materiałowa 1 (2001) 46-49.

[7] W. Szkliniarz, Strukturalne aspekty wytwarzania stopów na osnowie faz międzymetalicznych z układu Ti-Al, Rudy i Metale Nieżelazne 9 (2002) 434-438.

[8] W. Szkliniarz, Doświadczenia w zakresie wytwarzania i przetwarzania stopów na osnowie fazy międzymetalicznej TiAl, Inżynieria Materiałowa 2 (2007) 47-53.

[9] J. P. Kuang, R. A. Harding, J. Campbell, Investigation into refractories as crucible and mould materials for melting and casting γ−TiAl alloys, Materials Science and Technology 16 (2000) 1007-1015.

[10] A. Kostov, B. Friedrich, Selection of crucible oxides in molten titanium and titanium aluminium alloys by thermo-chemistry calculations, Journal of Mining and Metallurgy 41B (2005) 113-125.

[11] Q. Jia, C. C. Cui, R. Yang, Intensified interfacial reactions between gamma titanium aluminide and CaO stabilized ZrO2, International Journal of Cast Metals Research 17 (2004) 23-27.

[12] D. Eylon, Information on http://www.tokyo.afosr.af.mil/OldReports/9418.html.

[13] K. Kim, T. W. Hong, H. Lee, Y. J. Kim, Melting and casting of titanium alloys, Journal of the Korean Foundrymen’s Society 19 (1999) 210-215.

[14] K. Myoung-Gyun, S. Si-Young, K. Young-Jig, Induction melting and casting process of Ti and TiAl alloys, Proceedings of the 65th World Foundry Congress, Gyeongju 2002, pp. 587-593.

[15] J. Barbosa, A. Duarte, C. S. Ribeiro, F. Viana, C. Monteiro, Zr bearing γ−TiAl induction melted, Key Engineering Materials 188 (2000) 45-54.


Grain refinement of Ti-48Al-2Cr-2Nb alloy by heat treatment method

Agnieszka Szkliniarz 1,a 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland

a [email protected]

Keywords: TiAl based alloys, grain refinement, cyclic heat treatment, discontinuous coarsening

Abstract. In this paper, the possibility of refining grain of Ti-48Al-2Cr-2Nb alloy in the processes

of multi-stage heat treatment consisted of initial heat treatment, cyclic heat treatment and under-

annealing was evaluated. Microstructural changes that take place during the particular heat

treatment procedures were also described. It was demonstrated that due to the application of

combined cyclic heat treatment and under-annealing almost 24-fold grain refinement in relation to

the state after homogenising could be obtained. Probable mechanisms of grain refinement in the

proposed heat treatment processes were also presented and influence of individual procedures of the

proposed treatment on selected properties of the investigated alloy was described.

Introduction

Low plasticity at room temperature and under plastic working conditions as well as susceptibility to

brittle cracking significantly limit the use of two-phase TiAl intermetallic alloys as construction

materials [1÷8]. Unfavourable influence of coarse-grain microstructure and favourable influence of

its refinement on a number of mechanical properties place the problem of finding solutions for the

effective grain refinement in the circle of fundamental issues for this group of alloys [9÷11].

For grain refinement of TiAl intermetallic alloys, the modification procedures, and above all the

plastic working and recrystallisation annealing procedures are used most often. The use of the first

type of procedures is limited by presence of brittle intermetallic phases in microstructure of

modified alloys, which form modifiers and alloying components, while the use of the latter type is

limited by low deformability of alloys [12]. Taking these factors into consideration, it seems

appropriate to become interested in heat treatment procedures as alternative methods for grain

refinement of TiAl intermetallic alloys.

For grain refinement of TiAl intermetallic alloys, the heat treatment consisting of combined

procedures of hardening from α single-phase area temperature and tempering in the upper

temperature range of α+γ two-phase area is used frequently. The result is improvement in strength

properties of alloys subject to heat treatment. According to Hu et al. [13÷15], the increase in

strength properties of alloys subject to heat treatment is caused by grain refinement as a result of

recrystallisation of highly defected massive phase γm. This fact is not confirmed by other authors

[16, 17] who associate the increase in strength properties with obtaining the dispersion two-phase

α2+γ microstructure, formed as a result of decomposition of highly defected phase γm. They think

that obtaining massive phase γm after hardening is the precondition of successful performance of

this treatment and it is only possible for alloys with aluminium content above 46.5 at.%. Due to

specific massive transformation mechanism α→γm, this is only possible for products with initial

fine grain. For these reasons, the suitability of this heat treatment for grain refinement in this group

of alloys seems to be disputable and limited [8].

In the years 2000-2002, Wang et al. [18÷21] published the results of research on the application

of cyclic heating up to the α+γ or α phase temperature range followed by fast cooling for two-phase

TiAl intermetallic alloy grain refinement. They associate the obtained grain refinement effect with

cyclic massive transformation α→γm during fast cooling and reverse transformation γm→α during

fast heating. The need to conduct the cyclic heat treatment under fast (impact) heating and cooling

(water, polymer) conditions limits the possibilities of its practical use only to products with small

cross-sections and simple shapes (the presented results were obtained on samples of 10×10×10


mm). Reaching the conditions of impact heating for products with bigger cross-sections is

unattainable with furnace heating. In turn, the use of fast cooling (water, polymer), necessary to

induce the massive transformation α→γm, results in frequent cracking of treated semi-products

during cooling followed by fast heating.

Papers [22, 23] show that for grain refinement of conventional two-phase titanium alloys the

combined procedures of cyclic heat treatment and long-term annealing at a temperature of two-

phase area can be used. It was determined that under conditions of cyclic heat treatment, conducted

at moderate heating and cooling rates and with the appropriate selection of temperature and time

parameters, the defect density for alloys increases as result of phase transformation α+β↔β. Under

conditions of further annealing, this results in significant grain refinement.

Taking into consideration large conformity between conventional two-phase titanium alloys with

microstructure α+β and two-phase TiAl intermetallic alloy with microstructure α2+γ [23], the

effective method for grain refinement of coarse-grained TiAl intermetallic alloys, which is the

application of combined cyclic heat treatment and under-annealing procedures, is presented in this

paper.

Materials and research methodology

A two-phase Ti–48Al–2Cr–2Nb (at.%) alloy with chemical composition as specified in table 1 was

used like the most representative one in the group of two-phase TiAl intermetallic alloys.

Table 1. Chemical composition of research alloy

Element content, at. %

Al Cr Nb O Ti

48.22 1.97 2.02 0.43 Balance

The alloy was melted in the Leybold-Heraeus IS-5/III vacuum induction furnace, using the ceramic

crucible with applied ZrSrO3 layer. The alloy was cast into the preheated graphite moulds as bars of

13 mm in diameter and 120 mm in length. The bars were homogenised at 1400°C for 1 h followed

by cooling in furnace. After that the alloy was subjected to cyclic heat treatment (Fig. 1.-stage 2). A

single heat treatment cycle was continuous resistance heating at the rate of 20 to 100°C/s up to

temperature of 1360 to 1440°C (above Tα temperature) followed by cooling at the rate of 10 to

60°C/s to temperature below the eutectoid transformation temperature. The number of cycles was

changed from 1 to 20. The heat treatment was carried out with specially designed equipment using

direct resistance heating and gas stream cooling through the system of nozzles directed immediately

at the heated sample. The isothermal annealing (Fig. 1.-stage 3) was conducted at 1250, 1300 and

1350°C (temperature of α + γ two-phase area) for 2÷24 hours with air cooling after soaking.

The specimens for optical and scanning electron microscope (SEM) observation were

electrochemical polished at approx. -30°C and 20 V in the solution of 5% of perchloric acid, 35%

n-butanol and 60 vol.% of methanol, and etched in Kroll’s reagent. Thin foils for transmission

electron microscope (TEM) were prepared with a twin-jet electropolishing technique using

a solution of 13% perchloric acid, 74% nbutyl alcohol and 13% methanol at -30°C (25 V).

Microstructural examinations were performed using a Olympus GX71 light microscopes, Hitachi

S-4200 scanning electron microscope, and JEM-100B transmission electron microscope.

Volume fraction and size of phase crystallites were determined based on investigations carried

out using the X-ray diffractometer JEOL JDX-7S. Size and shape of grains were assessed using the

image analysis software Aphelion 3.2. The compression tests were performed on an INSTRON

4483 testing machine.


Research results

As a result of combining the procedures of heat treatment consisting of:

• initial heat treatment (Stage 1) – isothermal holding within the temperature range of α single-

phase area for 1 h,

• cyclic heat treatment (Stage 2) – the most important stage of the treatment the single cycle of

which consisted of fast continuous heating up to the α single-phase area temperature combined

with directly following cooling,

• under-annealing (Stage 3) – isothermal holding at different temperatures within the range of

α+γ two-phase area for up to 24 h and air cooling after holding,

almost 24-fold grain refinement as compared to the state after homogenising was obtained (Fig. 1).

Fig. 1. Multi-stage heat treatment diagram

Changes in microstructure during the successive treatment stages and their probable mechanisms

are presented in the following subsections.

Stage 1 - Initial heat treatment

For removal of the effects of dendritic microsegregation in cast TiAl-based alloys, homogenising

within the α single-phase temperature range is usually used (Fig. 2a, b).

Ti-48Al-2Cr-2Nb alloy after homogenising is characterised by coarse grain with average

equivalent diameter of grain plane section of 1200±625 µm and two-phase lamellar microstructure

consisting of γ phase, which is predominating, and α2 phase (Fig. 3a). Lamellas usually occur in the

γ/α2 arrangement, less widely in the γ/γ arrangement. They form colonies with identical orientation

within grains and with different orientation in particular grains. Crystallographic orientations

between the individual alloy phases maintain the relations determined by Blackburn [24] for this

group of alloys (Fig. 3b).


Fig. 2. Macro- and microstructure of Ti-48Al-2Cr-2Nb alloy: after casting (a) and after

homogenising: 1400ºC/1 h/cooling with furnace (b)

Changes in microstructure that accompany the α+γ↔α transformations are presented in Figs. 4

and 5.

Fig. 4. Microstructure of Ti-48Al-2Cr-2Nb alloy cooled in water immediately after heating with the

rate of 0.08°C/s up to: 1300 (a), 1310 (b), 1320 (c) and 1330ºC (d)

Fig. 5. Microstructure of Ti-48Al-2Cr-2Nb alloy after annealing at 1400°C for 1 h and cooling with

the rate of 0.08°C/s up to: 1250 (a), 1200 (b), 1100 (c) and 1050ºC (d) and further in water

Fig. 3. Microstructure of Ti-48Al-2Cr-2Nb alloy after homogenising with visible γ/γ and γ/α2

interfaces (a) and diffraction from the middle of the area in figure (a) with solution (b)

abγγγγ

γγγγ

γγγγ

γγγγ

γγγγ

αααα2

αααα2

αααα2

αααα2

γγγγ

a b c d

a b c d


They consist in dissolution, preceded by fragmentation, of lamellar precipitations of γ phase

during heating (Fig. 4) and formation of the grid of γ phase precipitations at grain boundaries and its

growth towards the depth of grain until its entire space is filled with lamellar microstructure during

cooling (Fig. 5).

Slow heating up to the α phase temperature range and isothermal holding within this range in

combination with slow cooling result in reconstruction of the initial microstructure without visible

symptoms of grain refinement.

During heating with significantly higher rate (35 and 100°C/s), even after temperature of 1400°C

has been reached, the α+γ→α transformation does not take place completely, which is proven by

the presence of undissolved precipitations of γ phase in microstructure of water-cooled alloy

(Fig. 6a). This phase occurs in the form of lamellas and fragments of lamellas within the grains and

at grain boundaries. Thickness and volume fraction of undissolved lamellar precipitations increases

with the increase in heating rate.

Fig. 6. Microstructure of Ti-48Al-2Cr-2Nb alloy cooled in water (a) and with the rate of 10°C/s

(b, c) immediately after heating up to 1400°C with the rate of 35°C/s

The undissolved lamellar precipitations of γ phase or their fragments, which remain in the alloy

microstructure after heating up to 1400°C with the rate of 35°C/s (Fig. 6b, c), effectively inhibit the

growth of lamellar microstructure α+γ tending to fill the entire space of the “old” grain during

cooling. These can also be the areas where nucleation of new grains takes place. The size of new

grains is related to the orientation of lamellar microstructure, which is formed during cooling. If it is

inconsistent with the orientation of undissolved lamellar precipitations of γ phase, its growth is

quickly inhibited (Fig. 7). In this way, new grains 1 and 4 are formed (Fig. 7b). When the

orientation of new lamellas is consistent with the orientation of undissolved lamellar precipitations

of γ phase, then their growth may proceed without any hindrance. In this way, grain 2 (Fig. 7b) will

probably reconstruct its original shape. It is also possible that nucleation of new grains will take

place on undissolved lamellar precipitations of γ phase (grain 3 in Fig. 7b).

Fig. 7. Microstructure of Ti-48Al-2Cr-2Nb alloy after heating up to 1400°C with the rate of 35°C/s

and cooling with the rate of 10°C/s (a) and its diagram (b)

a b

a b

c


The effect of undissolved lamellar precipitations of γ phase on the formation of microstructure

during cooling shows the possibilities of Ti-48Al-2Cr-2Nb alloy grain refinement in the processes

of heat treatment consisting of fast heating and cooling procedures.

Stage 2 – Cyclic heat treatment

As a result of the application of cyclic heat treatment, more than 8-fold reduction in average grain

diameter of Ti-48Al-2Cr-2Nb alloy with coarse-grained lamellar microstructure was obtained

(Fig. 8). The analysis of the effect of cyclic heat treatment parameters on grain refinement is

presented in papers [24, 26]. It was determined that the highest refinement effect occurred after the

5th

cycle of treatment performed under the following conditions: upper cycle temperature – 1400°C,

heating rate – 35°C/s, cooling rate – 10°C/s. Further increasing the number of cycles has slight

effect on changes in size (Fig. 9a) and shape (Fig. 9b) of grain and relative surface area of its

boundaries (Fig. 9c).

Fig. 8. Effect of cyclic heat treatment on macro- (a) and microstructure (b) of Ti-48Al-2Cr-2Nb

alloy

Fig. 9. Effect of the number of cycles on size, shape and relative surface area of grain boundaries

of Ti-48Al-2Cr-2Nb alloy


The microstructure of alloy subject to cyclic heat treatment is two-phase lamellar mixture of

crystals from α2 and γ phases with lamella thickness depending on the applied cooling rate from

upper to lower cycle temperature (Fig. 10). The higher the cooling rate is, the lower the lamella

thickness becomes. Regardless of the applied cooling rate from upper to lower cycle temperature,

the thickness of lamellas is always smaller than their thickness in the initial microstructure.

Fig. 10. Effect of cooling rate on microstructure of Ti-48Al-2Cr-2Nb alloy after the 1st cycle of heat

treatment performed under the conditions: upper cycle temperature – 1400°C, heating rate – 35°C/s

and different cooling rates from upper and lower cycle temperature – 10°C/s (a) and 45°C/s (b),

respectively

The area where lamellar microstructure nucleates is boundaries of grains with undissolved γ

phase (Fig. 11a), boundaries of undissolved lamellas of this phase (Fig. 11b) or clusters of its fine

undissolved particles (Fig. 11c). Growth of lamellar microstructure within the individual grains

probably takes place according to the “terrace-ledge-kink” mechanism, typical of these alloys

(Fig. 11d).

Fig. 11. Microstructure of Ti-48Al-2Cr-2Nb alloy after the 1st cycle of heat treatment performed

under the conditions: upper cycle temperature – 1400°C, heating rate – 35°C/s, cooling rate – 10°C/s

a b

c d

a b


Number of cycles has the effect on change in volume fraction of phases, which include in the

alloy microstructure (Fig. 12). With increase in the number of cycles the volume fraction of α2

phase in microstructure decreases from approx. 19% for the state after homogenising to approx. 7%

after 10 treatment cycles. Significant changes in volume fraction of α2 phase take place within the

range up to 10 cycles.

Fig. 12. Changes in volume fraction of α2 phase in microstructure of Ti-48Al-2Cr-2Nb alloy after

cyclic heat treatment

Changes also concern to γ phase. The content of γ phase in microstructure of alloy subject to heat

treatment is predominating. It was found that with increase in the cycles number the size of γ phase

crystallites was decreasing (Fig. 13a), while the value of crystal lattice deformation of this phase

was increasing (Fig. 13b). The biggest changes in these parameters took place within the range up

to 5 cycles.

Fig. 13. Effect of the cycles number on crystallite sizes (a) and changes in lattice deformations of

γ phase of Ti-48Al-2Cr-2Nb alloy

The analysis of changes in microstructure of Ti-48Al-2Cr-2Nb alloy after heating and cooling

with different rates and after cyclic heat treatment with different number of cycles allows the

scheme of change in microstructure and probable mechanism responsible for grain refinement as

a result of this treatment to be presented.


Fig. 14. Scheme of changes in microstructure during cyclic heat treatment

During heating of Ti-48Al-2Cr-2Nb alloy with lamellar microstructure up to the upper

temperature of the first treatment cycle the α+γ→α transformation occurs. As a result, lamellar

precipitations of γ phase are dissolved and the process consists in moving the α/γ interface towards

γ phase (Fig. 14). After the upper cycle temperature has been reached, there remain numerous

undissolved lamellar precipitations of γ phase due to high heating rate and large thickness of initial

lamellas in the alloy microstructure, in particular at γ/γ interfaces,. During the immediately following cooling, the undissolved lamellar precipitations inhibit the growth in lamellar

microstructure formed as a result of the α→α+γ transformation. During the next cycles, after heating to the upper cycle temperature, the number of undissolved lamellas and their fragments decreases as compared to the state after the first cycle. It is caused by significantly higher thickness of lamellas in the initial microstructure of alloy after homogenising as compared to their thickness after the first and following cycles. In the presence of numerous fragments of undissolved lamellar

precipitations of γ phase, their inhibiting effect on growth of lamellar microstructure, which nucleates within the boundaries of “old” grains and on fragments of other lamellas, trying to fill the area of the “old” grain, is visible during cooling that follows the successive cycles (Fig. 14). During the next cycles, a unique state of equilibrium between the number of undissolved lamellar

precipitations of γ phase and the number of grains is established. Thus, grain size is not subject to any substantial changes, however grains become more equiaxial and homogenous with regard to their sizes and shapes (Fig. 9a, c).

The presented diagram of changes in microstructure (Fig. 14) during the cyclic heat treatment processes and the grain refinement mechanism indicate a decisive share of undissolved lamellar

precipitations of γ phase in the process of forming the final grain size using methods of this treatment. This is the reason, among other things, for so large influence of the upper cycle temperature and heating rate on the obtained grain refinement effect.

During heating to the upper cycle temperature, the processes of dissolving γ phase as a result of

the α+γ→α transformation take place. With increase in the upper cycle temperature, the progress of

the α+γ→α transformation is increasingly higher, but even at 1440°C, which is close to the melting point, it is not complete due to high heating rate. When the upper cycle temperature is too low, refinement does not occur at all or concern the areas nearby the boundaries of “old” grains only.

When the temperature is too high, the degree of the α+γ→α transformation is so high that there is

no sufficient number of undissolved lamellar precipitations of γ phase, which are the areas where nucleation of new grains takes place and which inhibit their growth. It results in bigger grain size. In addition, too high upper cycle temperature is danger because of partial melting.


The change in heating rate from lower to upper cycle temperature has a similar effect as the

upper cycle temperature value. The increase in the rate of heating to the same upper cycle

temperature causes that the number of undissolved lamellar precipitations of γ phase is increasingly

higher. Too low heating rate causes that the number of these precipitations is small, which affects

lower grain refinement. The highest refinement occurs at the heating rate of 35°C/s. Large number

of undissolved lamellar precipitations of γ phase, which occurs after heating with the rates of 50 and

100°C/s, causes that they more determine the orientation of the newly lamellar microstructure

formed during cooling. Thus, they more promote reproduction of the initial grain size in these areas

than inhibit growth of new grains with lamellar microstructure.

Within the examined range of cooling rates (10÷60°C/s), refinement concerns the entire grain

area at the lowest cooling rate only. High cooling rate causes that new grains with lamellar

microstructure, growing from the boundaries of “old” grains, will not be able to reach large sizes

and form a ring at their boundaries, with middle area of the grains remained unchanged. Reduction

in cooling rate causes that refinement covers a bigger and bigger area of the “old” grain. Out of all

the analysed parameters, the cooling rate from the upper to bottom cycle temperature has the

biggest influence on the grain refinement effect. In this case, beside the role of undissolved lamellar

precipitations of γ phase, the inhibiting influence of high cooling rate on the growth of new grains is

also visible in the final effect of grain refinement.

Stage 3 – Under-annealing

Immediately after cyclic heat treatment, as a part of the third stage of the proposed heat treatment

(Fig. 1), the procedures of under-annealing conducted at the α+γ two-phase area temperature were

applied. The microstructure changes of Ti-48Al-2Cr-2Nb alloy subjected to under-annealing

indicate the occurrence of the “discontinuous coarsening” process already after 1÷2 h. The effects

of discontinuous coarsening usually occur in such alloys only after long-term annealing or operation

for 1000 hours or more [27÷29]. It follows, that cyclic heat treatment significantly accelerates the

course of the discontinuous coarsening process that occurs during under-annealing performed at the

α+γ two-phase area temperature, and this fact should be associated with microstructure instability.

The time after which the effect of discontinuous coarsening process includes the whole area of

grain remains in close relation with under-annealing temperature. At 1200°C, it takes place after

24 h (Fig. 15), at 1300°C – after 16 h (Fig. 15), and at 1350°C – already after 8 h (Fig. 15). The

discontinuous coarsening process is accompanied by degradation of lamellar microstructure. In

annealing performed at 1200°C the degradation occurs simultaneously with the discontinuous

coarsening process, while in annealing at 1300 and 1350°C the degradation of lamellar

microstructure is activated only when the discontinuous coarsening process includes the whole area

of grain (Fig. 15). The processes of lamellar microstructure degradation consist in fragmentation

and coagulation of lamella fragments and begin at boundaries of the newly formed grains.

The analysis of changes in microstructure of Ti-48Al-2Cr-2Nb alloy subject to under-annealing

at the α+γ two-phase area temperature allowed the diagram of probable course of the discontinuous

coarsening process resulting in grain refinement to be presented (Fig. 16). In accordance with

microstructure observations, the discontinuous coarsening process begins at the “old” grain

boundary and moves towards the depth of grain. It may also begin inside the grain and include only

a fragment of lamellar microstructure of the “old” grain. The process is diffusive in nature and

accompanied by relocation of components before its front. Its driving force, according to Livingston

and Cahn’s theory [30], is the natural tendency of every system to reduce its free energy by

reduction in total area of interfaces. The process takes place until its fronts, which move

independently from each other, contact together including the whole grain. This determines the final

shape and size of new grains, and thus the obtained refinement effect. New grains, which are the

result of the discontinuous coarsening process, are characterised by significantly higher thickness of

lamellas and orientation different from the orientation of lamellar microstructure of the “old” grain.


1200°C 1300°C 1350°C

4h

8h

16

h

24

h

Fig. 15. Effect of temperature and time of under-annealing on microstructure of Ti-48Al-2Cr-2Nb

alloy previously subject to cyclic heat treatment

Fig. 16. Diagram of changes in microstructure during under-annealing of alloy previously

subjected to cyclic heat treatment

The changes of microstructure of Ti-48Al-2Cr-2Nb alloy after proposed heat treatment are in

correlation in changes of properties. The application of cyclic heat treatment results in increase in

strength as compared to the state before treatment by approx. 130 MPa (Tab. 1). Grain refinement


that accompanies the cyclic heat treatment also results in increase in plasticity from 3.6 to 5.8%

with reference to the state after homogenising. Under-annealing conducted at 1300°C for 16 h of

alloy previously subject to cyclic heat treatment leads to further grain size refinement what results

in further improvement in plastic properties (Tab. 1). As a result of this the compression strength is

reduced by 115 MPa and the peak flow stress decrease by approx. 40 MPa compared to state after

cyclic heat treatment (Tab. 2).

Table 2. Effect of heat treatment on mechanical properties of Ti-48Al-2Cr-2Nb alloy –

compression test

State Rc

[MPa]

Rc0,2

[MPa]

Ac

[%] HV1

As cast 1290 1055 1.7 -

Homogenization: 1400ºC/1 h/furnace 1350 1085 3.6 250

Cyclic heat treatment (5 cycles) 1480 1140 5.8 350

Cyclic heat treatment + annealing 1365 1020 6.5 330

Table 3. Effect of heat treatment of alloy register in hot compression test

Parameter Heat treatment

Homogenization Cyclic heat treatment Cyclic heat treatment + annealing

σpm, [MPa] 676 574 531

εp 0.148 0.082 0.077

Summary

Multi-stage heat treatment consisting of homogenising, cyclic heat treatment and under-annealing

allows the efficient grain refinement process of TiAl intermetallic alloys.

The phase transformations that occur during fast heating and cooling from lower to upper

temperature of thermal cycle are responsible for the grain refinement effect in the cyclic heat

treatment processes. The main role in the refinement process is played by undissolved lamellar

precipitations of γ phase, remaining in the microstructure after fast heating to the α single-phase

area temperature. During cooling the γ phase precipitations are places of new grains with lamellar

morphology nucleation and inhibit their growth.

As a result of application of cyclic heat treatment, more than 8-fold reduction in average grain

diameter of the investigated alloy was obtained.

The effect of the discontinuous coarsening process, which occurs during under-annealing

performed immediately after cyclic heat treatment, are new grains with higher lamella thickness,

nucleating at the boundaries of “old” grains, or less widely inside them, and growing until they

contact together and include the whole area of grain. This results in another, 3-fold reduction in the

average grain diameter as compared to the state after cyclic heat treatment.

The grain refinement of TiAl intermetallic alloy by applying cyclic heat treatment is an efficient

method for improvement in mechanical and plastic properties as well as hardness of alloys. The

application of under-annealing after cyclic heat treatment results in further increase in plastic

properties. At the same time, it results in reduction in strength properties and hardness, which is

connected with the increase in lamella thickness in microstructure of the newly formed grains.

Proposed heat treatment processes can be successfully applied as a final heat treatment of cast or

initial heat treatment for forming ingots made of Ti-48Al-2Cr-2Nb alloy increased their plasticity.


References

[1]. H. Clemens, F. Appel, A. Bartels, Processing and application of engineering γ−TiAl based

alloys, in: Ti-2003 Science and Technology, WILEY-VCH Verlag GmbH & Co. KGaA,

Weinheim, 2004, pp. 2123-2137.

[2]. Titanium and Titanium Alloys, edited by Leyens C., Peters M., WILEY-VCH Verlag GmbH &

Co. KGaA, Weinheim, 2003.

[3]. I. J. Polmear, Recent developments in light alloys-overview, Mat. Trans. JIM 1 (1996) 12-31.

[4]. S. Naka, Advanced titanium based alloys, Solid State and Mat. Sci. (1996) 333-339.

[5]. S. Djanarthany, J.-C. Viala, J. Bouix, An overview of monolithic titanium aluminides based on

Ti3Al and TiAl, Mat. Chem. and Phys. 72 (2001) 301-319.

[6]. L. Changhui, X. Qiang, S. Yong-gian, Phase orientation relationships in the TiAl-TiAl2 region,

Mat. Sci. and Eng. A31 (2001) 227-236.

[7]. Y.-W. Kim, Ordered intermetallic alloys, part III: Gamma titanium aluminides, J. of Met. 46

(1994) 30-40.

[8]. Y.-W. Kim, D. M. Dimiduk, Progress in the understanding of gamma titanium aluminides, J. of

Met. 43 (1991) 40-47.

[9]. E. A. Loria, Quo vadis gamma titanium aluminide, Intermetallics 12 (2001) 997-1001.

[10]. J. C. Williams, J. Starke, A. Edgar, Progress in structural materials for aerospace systems,

Ac. Mat. 19 (2003) 5775-5799.

[11]. H. Clemens, H. Kestler, Processing and applications of intermetallic γ−TiAl-based alloys,

Adv. Eng. Mat. 9 (2000) 551-570.

[12]. S. L. Semiatin, V. Seetharaman, I. Weiss, Hot workability of titanium aluminide alloys – an

overview, Mat. Sci. and Eng. A243 (1998) 1-24.

[13]. D. Hu, P. A. Blenkinsop, M. H. Loretto, Alpha phase decomposition during continuous

cooling in Ti-48Al-2Cr-2Nb with and without boron addition, Proceedings of Ninth World

Conference on Titanium, Saint-Petersburg (1999) 290-297.

[14]. D. Srivastava, D. Hu, I. T. H. Chang, M. H. Loretto, The influence of thermal processing

route on the microstructure of some TiAl-based alloys, Intermetallics 7 (1999) 1107-1112.

[15]. D. Hu, R. R. Botten, Phase transformation in some TiAl-based alloy, Intermetallics 10

(2002) 701-715

[16]. W. J. Zhang, L. Francesconi, E. Evangelista, A novel heat treatment to develope very fine

lamellar structure in cast gamma-base TiAl alloys, Mat. Let. 27 (1996) 135-138.

[17]. C. E. Wen, C. Q. Chen, New microstructure-property approach: quenching / tempering

treatment in gamma TiAl alloy, Sc. Metal. et Mater. 8 (1997) 1283-1288.

[18]. Y. Wang, J. N. Wang, Q. Xia, J. Yang, Microstructure refinement of a TiAl alloy by beat

treatment, Mat. Sci. and Eng. A293 (2000) 102-106.

[19]. J. N. Wang, Q. Xia, Refining of coarse lamellar microstructure of TiAl alloys by rapid heat

treatment, Intermetallics 8 (2000) 545-548.

[20]. J. N. Wang, J. Yang, Q. Xia, Grain refinement of a TiAl alloy by heat treatment through

near gamma transformation, J. of Mat. Sci. 36 (2001) 4465-4468.


[21]. J. N. Wang, J. Yang, Q. Xia, J. Wang, On the grain size refinement of TiAl alloys by cyclic

heat treatment, Mat. Sci. and Eng. A329-331 (2002) 118-123.

[22]. W. Szkliniarz, G. Smołka, Analysis of volume effects of phase transformation in titanium

alloys, Journal of Materials Processing Technology, v. 53 (1995) 413-422.

[23]. W. Szkliniarz, Dwufazowe stopy tytanu-od klasycznych do intermetalików, Inżynieria

Materiałowa 6 (2003) 384-387.

[24]. M. J. Blackburn, The science, technology and application of titanium alloys, eds. R. I.

Jaffee, N. E. Promisel, Plenum, New York,1970.

[25]. A. Kościelna, W. Szkliniarz, Quantitative description of TiAl based alloy microstructure

after cyclic heat treatment, Inżynieria Materiałowa 4 (2008) 329-334.

[26]. A. Szkliniarz, W. Szkliniarz, Effect of cyclic heat treatment parameters on the grain

refinement of Ti-48Al-2Cr-2Nb alloy, Mat. Charact. 60 (2009) 1158-1162.

[27]. R. V. Ramanujan, Phase transformations in γ based titanium aluminides, Inter. Mat. Rev. 6

(2000) 217-239.

[28]. A. Denquin, S. Naka, Phase transformation mechanisms involved in two-phase TiAl-based

alloys - II. Discontinuous coarsening and massive-type transformation, Acta Mater. 44

(1996) 353-365.

[29]. R. Yamamoto, K. Mizoguchi, G. Wegmann, K. Maruyama, Effects of discontinuous

coarsening of lamellae on creep strength of fully lamellar TiAl alloys, Intermetallics 7-8

(1998) 699-702.

[30]. J. D. Livingston, J. W. Cahn, Discontinuous coarsening of aligned eutectoids, Acta Metal.

22 (1974) 495-503.


Characteristics of corrosion resistance of Ti-C alloys

Agnieszka Szkliniarz 1,a, Rafał Michalik 1,b 1 Silesian University of Technology, ul. Krasińskiego 8, 40-019 Katowice, Poland


Keywords: titanium alloys with carbon, corrosion test, corrosion resistance

Abstract. This paper presents the results of testing the corrosion resistance of pure Ti and Ti6Al4V

alloy improved by carbon addition at the level of 0.2 and 0.5 wt.%. The testing was carried out at

room temperature in HNO3 acid solution (40%) and HCl acid solution (5 and 10%). It has been

established that carbon addition affects the improvement in electrochemical corrosion resistance of

pure Ti and Ti6Al4V alloy in HNO3 solution, whereas the higher carbon content the better

corrosion resistance of Ti. For Ti6Al4V alloy the increase in corrosion resistance is caused by

carbon addition at the level of 0.2 wt.%. The result of the corrosion resistance of both pure Ti and

Ti6Al4V alloy with carbon in a solution of HCl indicates that the more detrimental is the solution of

lower concentration.

Introduction

The use of titanium and its alloys in the most demanding fields of technology arise from favourable

combination of strength and corrosion properties [1÷4]. The excellent corrosion resistance of

titanium and titanium alloys results from the formation of very stable, continuous, highly adherent,

and protective oxide films on metal surfaces, with thickness from 1.5 to 25 nm depending on

corrosion environment. The passive layer occurs immediately when metal surface is exposed to air

or humidity. It consists mostly of TiO2 oxide or TiO and Ti2O3 oxides distributed on the surface,

located mainly at the metal/oxygen interfaces [5]. Titanium shows excellent corrosion resistance in

many chemically diverse environments. It is completely resistant to: nitric acid, hypochlorous acid,

hydrogen sulphide, ammonia, seawater and many different chloride and sulphide solutions [5÷7].

The addition of elements such as Mo, Ni, Ta or Nb influences the increase in corrosion

resistance of titanium. Platinum metals in this regard show the highest efficiency. The Ti-0.15% Pd

alloy (Grade 7) is characterised by the highest corrosion resistance out of all Ti alloys.

Unfortunately, Pd addition results in two- or even three-fold increase in its price. Therefore cheaper

additions to allow the increase in both corrosion resistance and strength properties are sought after.

It is expected that such an alloy addition may be carbon considered so far as contamination in

titanium alloys [8]. The increase in strength properties caused by the presence of carbon takes place

at the cost of reduction in plasticity, impact resistance and susceptibility to cold forming to a level

acceptable only when carbon content in titanium does not exceed 0.5 wt% [9, 10].

The influence of carbon content on properties of titanium alloys is wide and characteristic for

individual alloy groups. Beside to the increase in strength properties, carbon addition may also

result in the increase in microstructure stability at elevated temperature, and even the increase in

plasticity at the appropriate ratio of carbon to oxygen content in alloy [11, 12]. A slight carbon

addition to pure titanium, at the level of 0.15 wt.%, results in 50% reduction in corrosion rate in the

environment of boiling 40% nitric acid solution [13]. 1.5% carbon addition affects the reduction in

titanium corrosion rate in the environment of boiling 3% hydrochloric acid solution (while 2% C

addition in the environment of boiling 5% HCl solution) to a level lower than for alloy with 0.15%

Pd addition. Carbon addition also has influence on the increase in stress corrosion resistance [13].

Taking into account these results in this paper presents the influence of carbon concentration on

electrochemical corrosion resistance of pure titanium and Ti6Al4V alloy in different corrosion

environments.


Material and Research methodology

For investigation, pure titanium and Ti6Al4V alloy with carbon content of 0.2 and 0.5 wt.% was

used. The alloys were made in vacuum induction furnace with the so-called “cold crucible”. Ingots

with diameter of 45 mm and length of approx. 250 mm were put to homogenising followed by hot

working. Rolled bars with diameter of 12 mm were put to recrystallisation annealing before

corrosion resistance testing.

Taking into account application of research titanium alloys for such elements as vessel or heat

excharger for tests selected strong aggressive and/or oxidizing acid. The corrosion resistance testing

was carried out in 40% HNO3 acid solution and 5 and 10% HCl acid solution. The DC

electrochemical measurements were taken in the conventional three-electrode system consisting of

the measuring cell and potentiostat Solartron 1285. The potentiodynamic testing of samples was

carried out in the range including cathodic and anodic potentials. Potential was changed within the

range Ecor(NEW) = -300 mV to 5000 mV at the rate of 10 mV/min. Temperature of solutions was

maintained at 21°C during the measurements. Test results were developed using CorrView 2

software. Corrosion current density was determined from the polarization resistance using Stern-

Geary equation. Polarization was conducted in a ranges not much different from the corrosion

potential for the observed linear dependence between current density and the sample potential. Test

pieces were grounded with abrasive papers.

The examination of alloy surfaces after corrosion tests was carried out on scanning microscopes

Hitachi S-4200 and S-3400N equipped with Thermo EDS detector.

Research results

The results of corrosion resistance testing of pure Ti and Ti6Al4V alloy with carbon addition are

presented in the form of sets of potentiodynamic curves (Fig. 1, 2).

I [A

/cm

2]

Ti- -- --->

<- ---- -Ti+0,2%C

<----- -Ti+0,5%C

0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0

10-7

10-6

10-5

10-4

10-3

I [A

/cm

2]

Ti6Al4V------>

<------Ti6Al4V+0, 2%C

<------Ti6Al4V+0,5%C

0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0

10-7

10-6

10-5

10-4

10-3

E (Volts)

E [V] E [V]

Ti Ti6Al4V

Fig. 1. Set of potentiodynamic curves for Ti and Ti6Al4V with different C content, tested in 40%

HNO3 solution


5% HCl 10% HCl

I [A

/cm

2]

Ti -- -- - ->

<--- - --Ti+0,2%C

Ti+0,5%C-- -- -->

-0.5 0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0

10-9

10-8

10-7

10-6

10-5

10-4

10-3

E (Volts)

I [A

/cm

2]

Ti-- - -- -><--- - --Ti+0, 2%C

Ti+0, 5%C---- - ->

-1.0 -0.5 0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0

10-8

10-7

10-6

10-5

10-4

10-3

E (Volts)

E [V] E [V]

Ti

I [A

/cm

2]

Ti6Al4V--- - --> <- -- - --Ti6Al4V+0, 2%C

Ti6Al4V+0,5%C--- -- ->

0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0

10-9

10-8

10-7

10-6

10-5

10-4

10-3

E (Volts)

I [A

/cm

2]

<---- --Ti6Al4V

<-- --- -Ti6Al4V+0,5%C

Ti6Al4V+0,2%C--- --->

-1.0 -0.5 0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0

10-8

10-7

10-6

10-5

10-4

10-3

E (Volts)

E [V] E [V]

Ti6Al4V

Fig. 2. Set of potentiodynamic curves for Ti and Ti6Al4V alloy with different C content, tested in

5% and 10% HCl solution

The results of corrosion resistance testing of Ti alloys with carbon content in 40% HNO3

solution indicate that under stationary conditions the highest corrosion resistance was shown by

titanium containing 0.5% C for which the lowest value of corrosion current density

(1.9×10-6

[A/cm2]) (Tab. 1) and the highest value of polarisation resistance (13 808 [Ohm/cm²])

were recorded (Fig. 3a). The passive range was observed on potentiodynamic curves (Fig. 1).

Table 1. Results of potentiodynamic tests of Ti alloys with C in different corrosion environments

Alloy

Corrosion environment

40% HNO3 5% HCl 10% HCl

Ecorr

[mV] Icorr ×10

-6

[A/cm2]

Ecorr

[mV] Icorr ×10

-7

[A/cm2]

Ecorr

[mV] Icorr ×10

-7

[A/cm2]

Ti 806 2.50 1090 0.32 -296 0.31

Ti+0,2% C 660 3.46 -111 1.87 305 0.28

Ti+0,5% C 760 1.90 314 2.38 -375 1.00

Ti6Al4V 560 1.40 342 0.17 352 0.87

Ti6Al4V+0,2% C 850 1.30 425 0.31 -372 0.84

Ti6Al4V+0,5% C 830 3.20 375 1.31 344 0.43

where Ecorr – corrosion potential, Icorr – corrosion current density


As it can be concluded from the curves (Fig. 1), carbon addition to pure titanium had a

favourable effect on electrochemical corrosion: for titanium with carbon addition the significantly

wider passive range was observed – the wider, the higher carbon content was. In the passive range,

alloys with carbon content were characterised by significantly lower passive current density than

that of pure Ti. It was 2.9 and 3.8×10-6

[A/cm²] for Ti containing 0.2 and 0.5 wt% C, respectively,

as compared to the value of 9.6×10-6

[A/cm²] characteristic of pure Ti. The peak on the

potentiodynamic curves for Ti alloys with carbon is the higher, the higher carbon content in Ti is.

Probably, this peak will be connected with dissolution of carbon precipitations existing in the

microstructure. It should also be noticed that higher carbon content in Ti results in moving the peak

in anodic direction.

a) b)

Fig. 3. Effect of C content and corrosion environment on polarisation resistance of tested Ti alloys

Table. 2. Results of potentiodynamic tests of Ti alloys with C in different corrosion environments

Alloy

Corrosion environment

40% HNO3 5% HCl 10% HCl

Ep-Ep-k

[mV] Ip ×10

-6

[A/cm2]

Ep-Ep-k

[mV] Ip ×10

-6

[A/cm2]

Ep-Ep-k

[mV] Ip ×10

-6

[A/cm2]

Ti 1410-2600 9.6 1930-2810 3.0 112-2550 1.3-4.5

Ti+0,2% C 780-1650 2.9 390-2280 1.3 903-2800 6.4-50.4

Ti+0,5% C 875-1980 3.5 640-2180 2.4-4.4

415-2790 10.0-27.3

Ti6Al4V 825-2070 8.5 610-1580 2.5 500-2700 1.5-15.5

Ti6Al4V+0,2% C 1345-1980 3.3 906-2530 0.5-1.9 275-2170 3.6-14.5

Ti6Al4V+0,5% C 1710-2030 5.7 540-2610 0.4-5.2 490-2670 2.2-6.7

where Ep-Ep-k – respectively the beginning and the end of the passive range, Ip – passive current

density

After corrosion tests in 40% HNO3 solution, no distinct effects that could indicate unfavourable

impact of carbon addition on corrosion resistance in this environment were observed on the surface

of tested materials (Fig. 4). The scratches visible, after the tests, on the surface of tested alloys,

occurred as a result of preparation of sample surfaces for testing, indicate very slow corrosion

process. Lack of any pits on the surface of pure Ti shows that corrosion was uniform.

For Ti6Al4V, alloys with carbon content were characterised by higher corrosion potential

(Tab. 1). Taking into consideration the value of corrosion current density, under stationary

conditions, the alloy containing 0.2% C is characterised by better corrosion resistance as compared

to the alloy with higher C content. The increase in carbon content in the tested alloy results in

increase in corrosion current and reduction in polarisation resistance from 19 848 Ohm/cm², which

is characteristic of alloy with 0.2% C content, to 8 172 Ohm/cm² (Tab. 1, Fig. 3). Under such

conditions, carbon addition results in decrease in corrosion resistance. However, it should be

noticed that stationary conditions are not the ones that would correspond to the real corrosion

conditions under which the tested alloy is exposed to the impact of local corrosion microcells. In the


passive range, the carbon-containing alloys were characterised by lower value of passive current

(3.3 and 5.7×10-6

A/cm² for alloy containing 0.2 and 0.5 wt.% C, respectively) as compared to the

alloy with no carbon content (8.5×10-6

A/cm²), which indicates the favourable effect of this element

on reduction in the corrosion rate (Tab. 2). However, carbon addition affects the reduction in

passive range (Tab. 2). This range is widest for the alloy containing no carbon. By analysing the

obtained results, it can be stated that carbon addition in the amount of 0.2 wt.% in Ti6Al4V alloy

results in increase in corrosion resistance. With its contents of approx. 0.5 wt.%, the effect on

corrosion resistance can be considered as neutral.

Ti Ti+0,2% C Ti+0,5% C

Ti6Al4V Ti6Al4V+0,2% C Ti6Al4V+0,5% C

Fig. 4. Surface appearance of alloys after corrosion resistance testing in 40% HNO3 solution

The observation of Ti6Al4V alloy surfaces after corrosion resistance testing in the environment

of 40% HNO3 solution indicates the local nature of corrosion (Fig. 4). As a result of the impact of

the applied corrosion environment, the surface of tested samples was etched.

Based on the results of corrosion resistance testing in 5% HCl solution, it was found that under

stationary conditions the increase in carbon content resulted in increase in corrosion current density

(Tab. 1) and decrease in polarisation resistance (Fig. 3b) in both pure Ti and Ti6Al4V alloy. It also

resulted in the reduction in corrosion potential, while significantly greater differences were

observed for pure Ti (Tab. 1). Carbon addition in both the tested materials affects the increase in the

passive range (Tab. 2). With carbon content of 0.2 wt.% in pure Ti, the value of passive current

density is getting reduced. For alloy Ti6Al4V with 0.2 and 0.5 wt.% C and Ti-0.2 wt% C, taking

into consideration the increase in the current density in this range (Fig. 2, Tab. 2), it can be said

about the pseudo-passive range. Taking into account low corrosion current density and high

polarisation resistance as well as low passive (pseudo-passive) current density with large width of

the passive range, it can be considered that 0.2 wt.% C additions to Ti6Al4V alloy has no influence

on deterioration of its corrosion resistance.

The impact of 5% HCl solution on the surface of pure Ti and Ti6Al4V alloy results only in its

insignificant dissolution (Fig. 5). The scratches visible on the surface indicate very slow and

uniform course of the corrosion processes. For pure Ti with carbon content the effects of corrosion

occur only locally on the surface in the form of fine pits. In particular, it is visible at higher contents

of carbon. For Ti -0.2 wt.% C, these are single areas.


On the surface of Ti6Al4V alloy with different carbon contents, craters characteristic of pitting

corrosion were observed after corrosion resistance testing carried out in 5% HCl solution. The way

of distribution of these craters indicates that these are probably areas occurred as a result of falling

out of carbides. Therefore their surface fraction increases with increase the carbon content in alloy

(Fig. 5). In addition, with higher carbon content in alloy the depth of these craters is higher, which

indicates more intensive corrosion. On the other hand, the scratches, which are still visible on the

surface of tested material after corrosion tests, show that corrosion takes place very slowly. In this

case, corrosion is of local nature. Taking into account very high value of breakthrough potential it

can be considered that the probability of the occurrence of this type of pits on the material surface

under real conditions is very low.

Ti Ti+0,2% C Ti+0,5% C


Fig. 5. Surface appearance of alloys after corrosion resistance testing in 5% HCl solution

The results of corrosion resistance testing in 10% HCl solution show that carbon addition to pure

Ti at the level of 0.2 wt.% results in slight decrease in corrosion current density (from 0.31×10-7

A/cm², which is characteristic for the initial state, to 0.28×10-7

A/cm²), insignificant increase in

polarisation resistance and increase in corrosion potential (from –296 mV, which is characteristic

for the initial state, to 305 mV) (Tab. 1). However, the passive range decreases significantly

(Tab. 2) and the value of passive current density increases as compared to pure Ti. Attention should

also be paid to the very high value of critical passivation potential (3 310 mV). Corrosion traces are

visible on the surface of Ti-0.5 wt.% C alloy only (Fig. 6). At lower carbon contents and in pure

titanium, corrosion takes place slowly and uniformly.

The corrosion resistance testing of pure and C-containing Ti6Al4V alloy in 10% HCl solution

indicates that at lower carbon content under stationary conditions the value of corrosion potential

decreases (from 352 mV, which is characteristic for the initial state, to –372 mV), while the


parameters such as corrosion current density and polarisation resistance remain on the same level

(Tab. 1). Under real conditions, the width of passivation range and the value of current density in

the passive (pseudo-passive) range remain unchanged.

Ti Ti+0,2% C Ti+0,5% C


Fig. 6. Surface appearance of alloys after corrosion resistance testing in 10% HCl solution

The changes observed on the surface of pure and C-containing Ti6Al4V alloy after the corrosion

resistance testing in 10% HCl solution (Fig. 6) are adequate with the changes after the impact of

5% HCl solution. Taking into consideration the obtained results, it can be stated that more

detrimental is the impact of HCl solution with lower concentration.

Summary

Based on the corrosion resistance testing in 40% HNO3 solution, it can be stated that carbon

addition to pure Ti has influence on the increase in corrosion resistance – the higher, the higher

carbon content was. For Ti6Al4V alloy the increase in corrosion resistance is observed at 0.2%

content of carbon. In 10% HCl solution, carbon addition to technical pure titanium results in

decrease in its electrochemical corrosion resistance and for Ti6Al4V alloy in 10% solution HCl –

the increase in corrosion resistance was observed at 0.5 wt.% C. However, there are pits on the

surface of this alloy, which probably occur at the points remained after dissolved carbides. In

5% HCl solution, under stationary conditions, technical pure titanium was characterised by lower

resistance to electrochemical corrosion than that in 10% HCl solution. A slight increase in corrosion

resistance was observed here for titanium with 0.2 wt.% C. In case of Ti6Al4V alloy in 5% HCl

solution the increase in corrosion resistance was observed for alloy with 0.2 wt.% C.

Acknowledgment

This scientific work is partially financed from the budget funds for science in the years 2008-2011

as the research project no NR15-0017-04.


References

[1] C. Leyens, M. Peters: Titanium and Titanium Alloys. WILEY-VCH Verlag GmbH & Co.

KGaA, Weinheim, 2003.

[2] I. V. Gorynin, N. F. Anoshkin: State of Production and Application of Titanium in the CIS.

Ti-2003 Science and Technology, WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim,

2004.

[3] A. Bylica, J. Sieniawski, Tytan i jego stopy, PWN, Warszawa, 1985.

[4] R. Melechow, K. Tubielewicz, W. Błaszczyk, Tytan i jego stopy: gatunki, właściwości,

zastosowanie, technologia obróbki, degradacja, Wyd. Pol. Częst., Częstochowa, 2004.

[5] R. W. Schulz, D. E. Thomas, Corrosion of titanium and titanium alloys, ASM Handbook,

Vol. 13, Corrosion, ASM International, Materials Park, OH 44073-0002, USA, (1987)

669-706.

[6] Materials Information Service – The Selection and Use of Titanium, A Design Guide,

Information on http://www.azom.com.

[7] J. A. Mountford, Titanium – properties, advantages and applications solving the corrosion

problems in marine service, Corrosion 2002, paper 02170.

[8] H. R. Ogden, R. I. Jaffee, The effects of carbon, oxygen, and nitrogen on the mechanical

properties of titanium and titanium alloys, Titanium Metallurgical Laboratory Report No. 20,

Ohio (1955) 1-101.

[9] B. A. Kolachev, A. V. Malikom, V. I. Sedov, Effect of carbon on structure and plasticity of

beta titanium alloys, Metall. i Term. Obrab. Met. 3 (1975) 42-43.

[10] Y. G. Li, P. A. Blenkinsop, M. H. Loretto, D. Rugg, W. Voice, Effect of carbon and oxygen

on microstructure and mechanical properties of Ti-25V-15Cr-2Al (wt%) alloys. Acta Mater.

10 (1999) 2889-2905.

[11] Z. Q. Chen, Y. G. Li, D. Hu, M. H. Loretto, X. Wu, Effect of carbon additions on the

microstructure and mechanical properties of Ti-15-3, Jo. of Mat. Sci. & Tech. 20 (2004)

343-349.

[12] M. Chu, L. P. Jones, X. Wu, Effect of carbon on microstructure and mechanical properties of

a eutectoid β titanium alloy, Jo. of Mat. Eng. and Perfor. 14 (2005) 735-740.

[13] S. J. Grauman, S. P. Fox, S. L. Nyakana, U.S. Patent, WO 2007/035422 A2.


Effect of a high-temperature hydrogen treatment on a microstructure and surface fracture in titanium Ti-6Al-4V Alloy

Maria Sozańska1,a


a [email protected]

Keywords: Ti-6Al-4V titanium alloy, high-temperature hydrogen treatment, hardness, fracture.

Abstract. Influence of hydrogen on the structure of titanium alloys is a complex phenomenon,

depending on the circumstances, may be negative or positive [1,2]. The presence of hydrogen in

titanium alloys usually results in degradation of their microstructure and properties, as well promote

some undesirable effects such as hydrogen corrosion and hydrogen embrittlement [3]. Positive

nature of the effects of hydrogen on the properties of titanium alloys is manifested in the high

temperature hydrogen treatment (HTM - Hydrogen Treatment of Materials), where hydrogen is

temporary alloying component [4-9]. This is possible because of the high values of diffusion

coefficients can be easily introduced into the titanium and it just as easily removed. Titanium and its

alloys show the absorbability of almost 60 at. % of hydrogen at 600°C. The limit hydrogen of

solubility in Tiα is very low and does not exceed 0.05 at. % at room temperature. The limit

hydrogen of solubility in Tiβ is much higher and its maximum value is 48 at. %. Since the beginning

of the titanium industry, a great deal of attention has been paid to control the hydrogen content at

titanium products – above 0.2 ppm. The paper presents the results of the possibilities of hydrogen

using as a temporary alloying element in Ti-6Al-4V alloy. Treatment of hydrogen alloy consisted of

three stages: hydrogenation in hydrogen gas atmosphere at 650 °C, a cyclic hydrogen-treatment (3

cycles 650 °C to 250 °C) and a dehydrogenation in vacuum (550 °C). It was shown that hydrogen

affects appreciably changes the microstructure of surface layer of the tested titanium alloy. The aim

of this study is thus to determine the effect of hydrogen on the two-phase microstructure, hardness,

and surface fracture of the titanium alloy Ti-6Al-4V due to high-temperature hydrogen treatment.

Introduction

The desirable impact of hydrogen on the properties of titanium alloys is manifested in high-

temperature hydrogen treatment (HTM: hydrogen treatment of materials), where hydrogen acts as a

temporary alloying element. In titanium alloys, hydrogen most often appears as hydrides. The most

stable is δ hydride (TiHx). Moreover, hydrogen alloying destabilizes the low-temperature α phase

and stabilises relatively ductile high-temperature β phase in two-phase Ti-6Al-4V alloy. By creating

interstitial type of solutions with individual allotropic forms of titanium, hydrogen significantly

changes the lattice parameters and specific volume of α and β phases, in particular of high-

temperature β phase for which it acts as the stabiliser. This results in significant internal-work

hardening during the α+β↔β transformations that proceeds under the heating and cooling

conditions. The presence of hydrogen also intensifies the eutectoid transformation process, as a

result of which the hydrogen-containing βΗ phase disintegrates into a mixture of α phase and δ

titanium hydride. The specific volume of titanium hydride is 13% to 17% higher compared with the

specific volume of the α phase, which causes large stress in the crystal lattice of this phase and may

result in its local plastic deformation [6-9]. The combination of cyclic heat treatment operations,

conducted at moderate heating and cooling rates, and recrystallization annealing within the range of

temperature of two-phase area may result grain size reduction at titanium alloys [5÷11]

The aim of this study is thus to determine the effect of hydrogen on the two-phase

microstructure, hardness, and surface fracture of the titanium alloy Ti-6Al-4V due to high-

temperature hydrogen treatment.


Material and experimental procedure

The study was of cylindrical specimens with a diameter of 16mm and a height of 12mm made of

titanium alloy Ti-6Al-4V content the following chemical composition (% mass.): Al-6.2, V-4, 3,

Fe-0.3, and Ti-the rest. Annealing was performed at 1100 °C for 1 hour, subsequent cooling of the

furnace.

Hydrogen treatment consists of the following three consecutive stages (Fig. 1):

1. hydrogenation (650°C, 1/2 h) – to obtain the microstructure that consists of α phase,

hydrogen-rich βH phase and δ titanium hydride,

2. cyclic heat treatment (650°C, 1/2 h, hydrogen, 3 cycles): low temperature heat treatment

using eutectoid transformation to obtain the appropriate microstructure defecting level as a

result of an internal work hardening caused by diversity of specific volumes of constitutive

phases occurring in the eutectoid that is repeated several times,

3. dehydrogenation (550°C, vacuum) – to remove hydrogen and, first of all, to force the

recrystallization the material having a microstructure defected due to cyclic heat treatment

that result in grain refinement.

Fig. 1. Schematic of high hydrogen treatment of titanium Ti-6Al-4V alloy.

The specimen-surfaces were etched in a solution of 50 cm3 of glycerol, 25 cm

3 HNO3, 30 cm

3

H2O, and 1 cm3 HF for metallographic examination.

A macrostructural examination of the specimen surfaces was performed using a stereoscopic

microscope (Olympus SZX-9). The microstructure and surface fracture was characterized by light

microscopy (LM, Olympus GX71) and scanning electron microscopy (SEM, Hitachi S-3400N).

Automated HV10 hardness tests were carried out on Struers Duramin A300 (HV10).

time



Figure 2 shows the specimens macrostructure of Ti-6Al-4V alloy before and after high

hydrogen treatment (HTM). Some difference of the dimension grain is visible.

(a)

(b)

Fig. 2. Macrostructure of titanium Ti-6Al-4V alloy (a) before HTM, (b) after HTM

The investigated Ti-6Al-4V alloy the in initial state, i.e. after full annealing within the β single-

phase area temperature was characterized by homogeneous two-phase microstructure with average

grain diameter of approx. 500 µm. The alloy microstructure is represented by alternately arranged

lamellar precipitations of α and β phases with different orientations in particular grains (Fig. 3).

(a)

(b)

Fig. 3. Microstructure of titanium Ti-6Al-4V alloy before HTM: (a) center specimen, (b) edge

specimen, LM.

Analysis of microstructural changes in the investigated alloys after high-temperature cyclic-

operation hydrogen treatment shows different microstructure changes in the center and at the edge

of the specimen. In the center, no significant changes in microstructure were observed (Fig. 4a, 5a).

Moreover, near the edge of the specimen, the lamellar precipitations of the α phases were divides

and the α phase relative volume was reduced (Fig.4b, 5b).

High-temperature hydrogen treatment led to fragmentation of lamellar precipitations of the

phase. The α phase dissolution process was limited to 50-150 µm from the edge of the specimen

(Fig. 6). The edge of the Ti-6Al-4V alloy specimen shows a lot of cracks after high-temperature

one- and three-cycle hydrogen treatment. The α phase also had significantly higher platelet phase

loss. At the same time, however, this was accompanied by a much greater number of microcracks.


(a)

(b)

Fig. 4. Microstructure of titanium Ti-6Al-4V alloy after HTM: (a) center specimen,

(b) edge specimen, LM.

(a)

(b)

Fig. 5. Microstructure of titanium Ti-6Al-4V alloy after HTM: (a) center specimen,

(b) edge specimen, SEM.

(a)

(b)

Fig. 6. Microstructure of titanium Ti-6Al-4V alloy after HTM: edge specimens, 3 cycle, SEM.


HV10 hardness results before and after high temperature hydrogen shown in Table 2. Analysis of

the hardness results indicates that significant differences exist between the sample before HTM and

after HTM. The results of the hardness of the hydrogen-treated samples indicate that the HV10

hardness in the layer is larger than in the middle of the samples.

Table 2. The results of HV10 hardness (before and after HTM)

.

Parameter

Treatment

Ti-6Al-4V alloy

Hardness

[HV10]

Center Edge

before HTM 314 ± 7.09 358 ± 29.34

after HTM,

3 cycles 329 ± 13.03 402 ± 27.88

Scanning electron microscopy (SEM) observation of the fracture surface for the samples heat

treated without hydrogen and with hydrogen treatment shows different type of fracture in the edge

part of specimens (Fig. 7). After hydrogen treatment, the fracture found of the brittle subsurface

zone in the edge part of specimens (Fig. 7c). This zone was approximately 100 µm. As in other

cases, the rest of the fracture was a ductile or quasi-ductile (Fig.7d).

(a)

(b)

(c)

(d)

Fig. 7. Surface fracture in the edge of Ti-6Al-4V section: (a) before HTM, (b) after HTM, SEM


Conclusions

1. The current state of knowledge and technological level allow for efficient forming of grain size

of two-phase titanium alloys only during hot or cold working processes in conjunction with

recrystallization annealing. The application of high-temperature hydrogen treatment may

significantly extend the possibilities for forming the microstructure of two-phase titanium

alloys.

2. The main effect of HTM is lamellar fragmentation of the alloy Ti-6Al-4V. The thickness of the

layer explicitly amended the HTM was up to 100 µm.

3. After the HTM revealed the presence of microcracks in the subsurface layer, this crack was

effect due to stress internal in the samples during the processing of hydrogen.

4. Fractal studies have confirmed that the samples in the initial state, without hydrogen, the

fracture morphology were a ductile or a quasi-ductile. However, after HTM, fracture

morphology, regardless of the parameters of a cyclic hydrogen treatment, showed a different

character in the layer samples than in the center. In the subsurface layer fracture morphology

was the fragile nature, while in the center ductile or quasi-ductile.

5. High-temperature hydrogen treatment with cyclic heat treatment is a very powerful method for

microstructure modification of the two-phase titanium alloys

Acknowledgments

This work was financed by budget funds for science in the years 2009-2012 as the research project

of the Ministry of Scientific Research and Information Technology No. N N507 462337.

References

[1] V. Tkachov: Mat. Sci. Vol 36 (2000), p. 481

[2] G. Solovioff and D. Eliezer: Scripta Met. Vol 40 (1999), p.1071

[3] V.A. Goltsov: J of Alloys and Comp Vol 293–295 (1999), p.844

[4] A. Zieliński: Niszczenie wodorowe metali nieżelaznych i ich stopów, Gdańskie Towarzystwo

Naukowe, Gdańsk, (1999) (in Polish)

[5] A. Takasaki, Y. Fyruya, K. Oima and Y. Taneda: J of Alloys and Compo Vol 224 (1999),

p.269

[6] D. Bhattacharyya, G.B. Viswanathan, R. Denkenberger, D. Furrer, Fraser, and L. Hamish:

Acta Materialia, Vol 51 (2003), p.4679

[7] O.N. Senkov: Mater Research Bulletin 36 (2001), p.1431

[8] O.A. Kaibyshev: J. of Materials Processing Technology Vol 117 (2001), p.300

[9] J. Nakahigashi and H. Yoshimura: J of Alloys and Comp Vol 330–332 (2002), p.384

[10] W. Simka, A. Iwaniak, G. Nawrat, A. Maciej, J. Michalska, K. Radwański and J. Gazdowicz,

Electrochimica Acta, 2009, 54, p.6983

[11] W. Szkliniarz: The work hardening cumulation effect during cyclic heat treatment of titanium

alloys, Inżynieria Materiałowa 3 (2002), p.96 (in Polish).


Diffusion brazing of titanium via copper layer

Maciej Różański1, a, Janusz Adamiec2, b 1 Institute of Welding, ul. Błogosławionego Czesława 16-18, 44-100 Gliwice, Poland



Keywords: titanium, diffusion brazing, vacuum, copper interlayer

Abstract

The paper presents the basic physico-chemical properties and brazability of titanium. The work also

discusses the principle and mechanisms of the formation of a diffusion-brazed joint and presents

results of metallographic and strength-related tests involving diffusion-brazed joints made of

technical titanium grade 2 via copper layer grade B-Cu100-1085. The paper also contains results of

structural examination conducted by means of light microscopy as well as results of shear strength

tests.

Introduction

Fast-developing aviation, automotive, power-generation and chemical industries cause an

increasing demand for new engineering materials which could resists such extreme operational

conditions as high operating temperature, considerable stresses or operation in fume-affected

environment. Highly desirable properties of such materials should include high strength, corrosion

resistance (also when exposed to aggressive fumes) and, last but not least, low density [5]. The

application of engineering materials combining all these features could ensure long and failure-free

life and reliability of components used in modern equipment and machinery operating under most

adverse conditions. Owing to such properties as low density (4.5⋅103 kg/m

3), high strength (tensile

strength of 500÷700 MPa) and excellent corrosion resistance, prospective structural materials

include titanium and its conventional alloys, which already find application in aviation, automotive,

power generation, chemical, petrochemical and food industries [1].

The industrial application of any structural materials requires the latter to be subjected to various

technological processes e.g. joining into a functional whole. As there is a number of difficulties

related to the welding of titanium and its commercial alloys, the brazing of the former appears to be

a highly promising joining method.

Being reactive, pure titanium belongs to materials which are difficult to braze [2, 3]. One of the

most convenient technical methods used for joining is diffusion welding. This method, combining

the qualities of brazing and diffusion welding, is usually referred to as ”a brazing process, in which

the mechanism of braze formation is based predominantly on the phenomenon of diffusion between

materials being joined and a brazing filler metal” or as ”a brazing process, in which the

phenomenon of diffusion is decisive for the chemical composition and physical properties of a

braze obtained by melting an added brazing filler metal or a brazing filler metal formed on the

contact point of elements being joined” [2-4]. The above definitions divide diffusion brazing into

two types. In the first type the filler metal is supplied from the outside and a chemical composition

of the liquid brazing filler metal is formed as a result of the mutual diffusion of components of the

filler metal and the base metal. It should be noted the melting point of both metals does not need to

have a melting point lower than the temperature at which a brazing process proceeds and that

intermetallic phases which are formed in the joint have melting points higher than the temperature

at which a brazing process takes place. An example of such a process is the diffusion brazing of

titanium with a copper wire (Fig. 1). The other type consists in brazing without a filler metal

supplied from the outside. A braze-forming liquid brazing filler metal is formed in the contact point

of materials being joined as their respective components undergo mutual diffusion. Such a


phenomenon occurs only in case of material systems whose components (or the materials

themselves) form phase equilibrium systems with a eutectic mixture or continuous solid solution

with a minimum on the liquidus curve. It is then that an alloy of eutectic composition or solid

solution composition with a minimum constitutes a brazing filler metal [2,3,4].

Fig.1. The Ti-Cu binary phase diagram with diffusion brazing temperature (broken line) [2]

Base and filler metals used in tests

A base metal used in tests was titanium grade 2 (max. content of impurities in % per weight:

0.1%C; 0.25%O; 0.03%N; 0.0125%H; 0.03%Fe) [6], out of which cylindrical test pieces were

made (dim. ∅ 20 x 15 mm);

An interlayer used in the process of diffusion brazing of titanium was made of 0.1mm-thick copper

wire grade B-Cu100-1085 [7].

Production of test joints

The elements of butt-brazed joints were placed freely and coaxially in the vertical position. In order

to increase the faying surface and diffusion length of the components of base metal and brazing

filler metal, prior to brazing the surfaces of the elements were subjected to grinding with abrasive

paper of the final designation of 800. Directly before brazing the elements were etched in aqueous

solutions of hydrofluoric acid and nitric acid. Shape-matched to a joint, the interlayers of copper

foil were degreased in acetone and placed between the elements to be joined.

All the samples were brazed in vacuum (range: 10-4÷10

-5 mbar) in S 16 TORVAC-made vacuum

furnace.

The brazing temperature and time were determined on grounds of available reference

publications and through the analysis of phase titanium-copper interaction on the basis of their

phase equilibrium systems [2,3,4]. Brazing was conducted at 900, 950, 1000 and 1030 °C, for 10,

20, 30 and 40 minutes (at each temperature). In each case heating up to the brazing temperature was

performed with 20-min-long isothermal holding at 700 °C in order to conduct desorption of gases

from the surface of elements being brazed.

The visual inspection of obtained joints revealed their good quality in case of the joints produced at

950, 1000 and 1030 °C. In turn, the test pieces brazed at 900 °C either completely failed to form a

joint or revealed only partial reacting of the brazing filler metal with the titanium base.


Structures of brazed joints of titanium grade 2

The test pieces for microscopic metallographic examination were subjected to grinding with

abrasive paper of gradation of 80, 320, 1000 and 2500 respectively and next to polishing by means

of polishing cloth with an addition of diamond and corundum polishing slurries of grain sizes of 3

and 0.05 µm respectively.

The microstructure of the brazed joints was revealed through etching of the samples in Buehler

etchant. The metallographic examination was carried out in the bright field using a Leica-

manufactured metallographic light microscope MeF4M.

The metallographic examination of brazed joints of titanium grade 2, made with a copper foil as

the brazing filler metal, revealed that in case of a short brazing time i.e. of 10 and 20 minutes, in the

central structure of joints there is a layer of unreacted interlayer metal, and on the boundary base

metal-braze one can observe a layer of phases of darker colour (Fig. 2 a and b). Along the boundary

of the said layer with the base metal and the braze it was possible to notice many cracks (Fig. 3 a

and b). In the joints made at a hold time of 30- and 40 minutes it diffused utterly to the base metal

(Fig. 2 c and d). In addition to that, in case of the joints brazed for 40 minutes it was possible to

observe a coarse-grained structure with acicular precipitates inside grains and an easily visible

boundary of crystallisation fronts (Fig. 2 d).

Fig.2. Microstructures of titanium (grade 2) joints diffusion brazed using interlayer of Cu filler

metal at 1000°C for 10 min (a), 20 min (b), 30 min (c), 40 min (d), etched with Buehler etchant

b) a)

c) d)


Fig.3. Crack on the boundary of brazed joints of titanium grade 2, made with copper brazing filler

metal B-Cu100-1085 (a) in the form of 0.1mm-thick foil, brazing temperature 1000°C, time 10 min

(a), 20 min (b), etched with Buehler etchant

Shear strength tests of brazed joints made of titanium

The strength properties of brazed cylindrical test pieces were determined by means of an Instron-

manufactured testing machine (model 4210) through shearing in special shackles designed in such a

manner that during shearing the samples were exposed to shearing forces only i.e. without bending.

The highest strength was revealed in case of the joints brazed for 30 minutes at a brazing

temperature of 950, 1000 and 1030°C and amounted to 245, 256 and 264 MPa respectively. The

shear strength of the test pieces brazed for 10, 20 and 40 minutes were contained in the range

83÷148 MPa.

The results of the static shear test of titanium test pieces brazed at various times and hold

temperature are presented in Fig. 4.

Fig.4. Impact of brazing time on shear strength of brazed joints of titanium grade 2, made with

copper brazing filler metal B-Cu100-1085 in the form of 0.1mm-thick foil, at brazing temperature

of 950, 1000 and 1030°C

a) b)

Sh

ear

str

en

gth

, M

Pa

Brazing time, min


Conclusions

1. The conducted material and technological tests enabled obtaining qualitatively proper

diffusion-brazed joints of titanium (grade 2) with the use of copper brazing filler metal

interlayers B-Cu100-1085; the applied brazing temperature amounted to 950 ÷ 1030 °C; the

brazing time was 10 ÷ 40 min.

2. Relatively high, repeatable shear strength values of the brazed joints of titanium (grade 2)

amounting to 245÷264 MPa were obtained at a brazing temperature of 950÷1030°C and a hold

time of 30 minutes.

3. A low shear strength of the joints brazed for 10 and 20 minutes was caused by cracks present

on the inter-phase boundary: braze – base metal as well as on the inter-phase boundary of the

individual layers of the phases present in the braze.

4. A probable reason for a decrease in the shear strength of the joints brazed for 40 minutes is

reduced coherence along the boundary of crystallisation fronts.

References

[1] L.A. Dobrzański: Basis of materials science. WNT, Warszawa 2002.

[2] Z. Mirski, M. Różański: Diffusion brazing of titanium and its alloys based on TiAl(γ) intermetallic compound, Materials Engineering, nr/no 2/2010, pp. 161-166.

[3] Z. Mirski, M. Różański: Vacuum brazing of TiAl48Cr2Nb2 casting alloys based on TiAl(γ)

intermetallic compound, Archives of Foundry Engineering, nr/no 1/2010, pp. 371-376.

[4] A. Winiowski: Impact of conditions and parameters of brazing stainless steel and titanium on

mechanical and structural properties of joints. Archives of Metallurgy and Metals, nr/no 4, pp.

593-607, 2007.

[5] J. Adamiec, T. Pfeifer, J. Rykała: Modern methods of aluminium alloys welding. Solid State

Phenomena, Vol 176, pp. 35-38.

[6] ASTM B 26579 Titanium and Titanium alloy Strip, Sheet and Plate.

[7] PN-EN ISO 17672 Brazing-Filler metal.


Keywords Index

A

Al3Ti 199

Aluminium Composites 81

Aluminium Refining 3, 13

Aluminum 29

Aluminum Alloy 37, 45, 75

Annealing 131

AZ61 Magnesium Alloy 101



B

Barbotage Process 13

C

Cast Aluminium Alloy 57

Castability 137

Casting Alloy 89

Ceramic Crucibles 211

CFD Simulations 89

CMT 45

Composite 57

Composite Material 67

Contact Resistance 89

Copper Interlayer 249

Corrosion Resistance 81, 235

Corrosion Test 235

Creep Properties 151

Creep Resistance 183

Creep Test 177

Cyclic Heat Treatment 221

D

Differential Scanning Calorimetry(DSC)

159, 189, 199

Diffusion Brazing 249

Discontinuous Coarsening 221

E

Elektron 21 Magnesium Alloy 123, 145

Elevated Temperature 131

Extrusion 169

F

Fine-Grained Microstructure 29, 37

Flow Stress 101

Fractography 109, 123

Friction Coefficient 67

G

Galvanic Corrosion 169

Glassy Carbon 67, 81

Grain Refinement 221

H

Hardness 137, 243

Heat Treatment 45, 177, 183

High-Temperature HydrogenTreatment

243

Hot Compression 101

Hybrid System 81

Hypereutectic Alloy 23

I

Infiltration 67

L

Low Energy Welding 45

M

Machinability 75

Magnesium Alloy 145, 151, 159,183

Magnesium Alloys AZ61 169

Magnesium Casting Alloys 109, 115, 123

Magnesium Composite 189, 199

MAXStrain 37

Mechanical Properties 37, 109, 123

Melting 211

Metal Matrix Composite (MMC) 75

Mg-5Al Magnesium Alloy 131

Mg-Al-Ca-Sr Alloy 151

Mg17Al12 199

Mg2Si 189


Microstructure 101, 109, 115,123, 131, 137,

151, 159

Microstructure Analysis 45

Minisamples 37

N

Nanoparticle Silica 189

Numerical Modelling 3

O

Overheating Degree 23

Oxidation 159

P

Physical Modelling 3, 13

Plastometric Tests 101

Porous Ceramics 57

Precipitation 131

Precursor 67

Pyrolysis 67

Q

QE22 and RZ5 MagnesiumCasting Alloys

137

QE22 Magnesium Alloy 109, 145

Quantitative Analysis 137

Quantitative Metallography 109, 115, 123

R

Repair Welding 177

RZ5 Magnesium Alloy 109

S

Severe Plastic Deformation (SPD) 29, 37

Silicon Carbide (SiC) 81

Simulation 137

Solid Lubricant 67

Solidification 89

Statistics 23

STEM 29, 37

Stepped Casting Test 145

Strain Hardening 37

Surface Geometry 75

T

TEM 29

Ti-6Al-4V Titanium Alloy 243

TiAl Based Alloy 211, 221

Titanium 249

Titanium Alloys with Carbon 235

Titanium Particles 199

V

Vacuum 249

W

WE43 Alloy 123

WE43 Magnesium Alloy 177

Weibull Distribution 23

Welding 183

X

X-Ray Tomography 189

Z

Zener-Hollomon Parameter 101

Authors Index

A

Adamiec, J. 45, 177, 183,249

B

Bednarczyk, I. 101

Boczkowska , A. 57

C

Chabera, P. 57

Chmiela, B. 151

Cwajna, J. 123, 137

D

Dolata, A.J. 57, 75, 81, 89

Dybowski, B. 109, 115, 123,137

Dyzia, M. 57, 75, 81, 89

H

Hadasik, E. 101, 169

J

Jarosz, R. 115, 137, 145

K

Kiełbus, A. 131, 137, 145

Kierzek, A. 177, 183

Kozera, R. 57

Kuc, D. 101

M

McDonald, S.A. 189

Merder, T. 3

Michalik, R. 235

Misiowiec, M. 199

Moskal, G. 189

Myalska, H. 189

Myalski, J. 67

O

Olszówka-Myalska, A. 189, 199

Oziębło, A. 57

P

Paśko, J. 109

Pawlicki, J. 29, 37

Pfeifer, T. 45

Piątkowski, J. 23, 159

Posmyk, A. 67

Przeliorz, R. 159, 199

Przondziono, J. 169

R

Rodak, K. 29, 37

Roskosz, S. 109, 115, 123

Różański, M. 249

Rykała, J. 45

Rzychoń, T. 151, 199

S

Saternus, M. 3, 13

Sozańska, M. 243

Stopyra, M. 145

Szkliniarz, A. 211, 221, 235

Szkliniarz, W. 211

T

Tkocz, M. 37

W

Walke, W. 81, 169

Wieczorek, J. 75

Withers, P.J. 189

Z

Zagórski, R. 89

light metals and their alloys ii : technology, microstructure and properties

Documents