Magnesium Alloys and Technologies

Magnesium Proceedings of the 6th International Conference Magnesium Alloys and Their Applications

Edited by K.U. Kainer

Vakat

Magnesium Proceedings of the 6th International Conference Magnesium Alloys and Their Applications Edited by K.U. Kainer

Further titles of interest: K.U. Kainer (Ed.) Magnesium Alloys and Their Applications (2003) ISBN 3-527-30570-X

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Magnesium Proceedings of the 6th International Conference Magnesium Alloys and Their Applications

Edited by K.U. Kainer

Editor: Prof. Dr. K. U. Kainer GKSS-Forschungszentrum Institut für Werkstoffforschung Max-Planck-Straße 21502 Geesthacht Germany

This book was carefully produced. Nevertheless, editor, authors, and publisher do not warrant the information contained therein to be free of erros. Readers are advised to keep in mind that statements, data, illustrations, procedural details or other items may inadvertently be inaccurate.

Library of Congress Card No.: Applied for British Library Cataloguing-in-Publication Data: A catalogue record for this book is available from the British Library Bibliographic information published by Die Deutsche Bibliothek Die Deutsche Bibliothek lists this publication in the Deutsche Nationalbibliografie; detailed bibliographic data is abailable in the Internet at . ISBN 3-527-30975-6

© 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim Printed on acid-free paper All rights reserved (including those of translation in other languages). No part of this book may be reproduced in any form – by photoprinting, microfilm, or any other means – nor transmitted or translated into machine language without written permission from the publishers. Registered names, trademarks, etc. used in this book, even when not specifically marked as such, are not to be considered unprotected by law. Composition, Printing and Bookbinding: Druckhaus “Thomas Müntzer” GmbH, Bad Langensalza Printed in the Federal Republic of Germany

Preface Since its launch in 1986, the International Conference on Magnesium Alloys and Their Applications has become the most reputable and excepted event as forum for current development in magnesium technology. The increase in interest on magnesium alloys for use in transportation industry and the dramatic changes in the magnesium industry in the last 5 years show consequences for the market development and influence the R&D activities. The 6th International Conference and Exhibition on Magnesium Alloys and Their Applications reflects this evolution process. The aim of the conference is to show the correlation of processing, microstructure and properties for magnesium materials. The complexity of interactions must be understood to use the full potential of magnesium materials for a particular application. The conference will cover the important aspects of magnesium research and technology including: Alloy Development Microstructure Evolution Casting Wrought Alloys Post Processing Recycling Simulation, Modelling Corrosion Surface Treatment Magnesium Matrix Composites Application Research Programmes The conference returns to Wolfsburg, the home of Volkswagen Group, where it saw the biggest success. The Call for Paper has attracted more than 230 papers from 28 countries. More than 190 contributions are published in the proceeding volume. They give an excellent overview on newest developments.

K. U. Kainer September 2003

Vakat

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Inhaltsverzeichnis Microstructural Design for Enhanced Elevated Temperature Properties in Sand-castable Magnesium Alloys C.J. Bettles and M.A. Gibson, CRC for Cast Metals Manufacturing (CAST), CSIRO Manufacturing & Infrastructure Technology, Private Bag 33, Clayton South MDC, Clayton, VIC. 3169, Australia .......................................................

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Thermodynamic Database for Mg-alloys: Progress and Application to Mg-Al-Ca-Si Alloy Development Rainer Schmid-Fetzer and Joachim Gröbner, University of Clausthal, Institute of Metallurgy, Clausthal-Zellerfeld, Germany .................................................. 12 Development of the New Creep Resistant Alloy AS31 Vadim V. Agalakov, JSC “AVISMA” , Berezniki, Russia ................................................ 18 Phase Formation, Precipitation and Strengthening Mechanisms in Mg-Zn-Sn Alloys S. Cohen, G.R. Goren-Muginstein, S. Avraham, and M. Bamberger, Department of Material Eng., Technion IIT, Haifa 3200 ISRAEL; G. Dehm, Max-Planck-Institut für Metallforschung, Stuttgart, Germany ....................... 25 Die Casting for High Performance – Focus on Alloy Development Dr. Per Bakke, Dr. Håkon Westengen, Hydro Aluminium a.s., Magnesium Competence Centre, Porsgrunn, Norway ........................................................................ 31 New Mg-Alloys for the Thixomolding™ Process Julio Aguilar, Martin Fehlbier, Andreas Bühring-Polaczek, Foundry Institute, Aachen University (RWTH), Germany............................................................................. 37 Structure and Morphology of Effective Obstacles in High Performance Mg – Rare Earth Base Alloys Bohumil Smola, Ivana Stulíková, Jitka Pelcová, Barry L. Mordike Faculty of Mathematics and Physics, Charles University, Prague 2, Czech Republic; Zentrum für Funktionswerkstoffe gGmbH, Clausthal-Zellerfeld; Institute of Materials Engineering and Technology, TU Clausthal, Clausthal-Zellerfeld, Germany ................ 43 Regularities in the Phase Diagrams and Behavior During Aging of the Ternary Mg-Base Alloys Containing Two Rare-Earth Metals of Different Subgroups Lazar Rokhlin, Tatyana Dobatkina, Nadezhda Nikitina, Baikov Institute of Metallurgy and Materials Science, Moscow, Russia ................................................... 49 Development of New Magnesium Alloys for Advanced Applications B. Bronfin, E. Aghion, S. Schumann, H. Friedrich, Dead Sea Magnesium Ltd, Beer Sheva, Israel; F. von Buch, Volkswagen AG, Wolfsburg, Germany ....................................................... 55

VIII Creep Resistance and Creep Kinetics of Mg-alloys W. Blum, Y. J. Li, X. H. Zeng, Universität Erlangen-Nürnberg, Inst. f. Werkstoffwissenschaften LS1, 91058 Erlangen, Germany; B. von Großmann, C. Haberling, H.-G. Haldenwanger, Audi AG, 85045 Ingolstadt, Germany......................................................................................................... 62 Creep Behaviour of MgZnRE Alloys Formed by Semi-Solid-Casting E. D. Morales, N. Hort, K. U. Kainer, GKSS Research Center Geesthacht GmbH, Geesthacht, Germany ....................................................................................................... 68 Creep Resistant Magnesium Alloys for Powertrain Applications Mihriban. O. Pekguleryuz, McGill, University, Department of Metals and Materials Engineering, University Street, Montreal, QC, Canada H3A 2B2; A. Arslan Kaya, Tubitak Marmara Research Center, P.O.B.21, Gebze, 41470 Turkey.. 74 Russian Ultralight Constructional Mg–Li Alloys. Their Structure, Properties, Manufacturing, Applications Fedor M. Elkin, Valentin G. Davydov, All-Russia Institute of Light Alloys, Moscow, Russia ............................................................................................................................... 94 The Effect of Grain Size on the Mechanical Properties of AM-SC1 G. Dunlop, Australian Magnesium Corporation, Milton, Australia; C.J. Bettles, J.R. Griffiths, K. Venkatesan, CRC for Cast Metals Manufacturing (CAST), CSIRO Manufacturing & Infrastructure Technology, Melbourne, Australia; Lihui Zheng and Ma Qian, CRC for Cast Metals Manufacturing (CAST), The University of Queensland, St Lucia, Australia.......................................................... 100 Thermal Properties of Mg-Li and Mg-Li-Al Alloys Alexandra Rudajevová, Pavel Lukácˇ, Department of Metal Physics, Charles University, Prague, Czech Republic; Stanislav Kúdela, Institute of Materials and Machine Mechanics SAS, Bratislava, Slovakia ............................................................................................................................ 106 Designing for Improved Creep Resistance in the RE – Free Magnesium Alloys Vladimir G. Tkachenko, International Center for Electronic materials Science and Applied Problems of Aerospace Technology, 3, Krzhyzhanovsky Str, Kyiv, 03680, Ukraine ....................................................................................................... 110 Ageing Characteristics and Creep Resistance of Mg-Y-Nd-Sc-Mn Alloy Ivana Stulíková, Bohumil Smola, Jitka Pelcová, Faculty of Mathematics and Physics, Charles University, Prague, Czech Republic; Ivana Stulíková, Bohumil Smola, Zentrum für Funktionswerkstoffe gGmbH, Clausthal-Zellerfeld, Germany; Barry L. Mordike, Institute of Material Science and Engineering, Technical University Clausthal, Clausthal-Zellerfeld, Germany ..................................................... 116 Deformation Behaviour of Mg-Li-Al Alloys at Room and Elevated Temperatures Zdeneˇk Drozd, Zuzanka Trojanová, Viera Gärtnerová, Faculty of Mathematics and Physics, Charles University, Prague, Czech Republic .................................................... 122

IX The Effect of Grain Size on the Bolt Load Retention Behaviour of AMC-SC1 C.J.Bettles, K.Venkatesan, CRC for Cast Metals Manufacturing (CAST), CSIRO Manufacturing & Infrastructure Technology, Melbourne, Australia; Lihui Zheng, Ma Qian, CRC for Cast Metals Manufacturing (CAST), The University of Queensland, St Lucia, Australia .................................................................................. 128 High Temperature Behaviour of a New Mg-Ca-Zn Alloy A. Zilberov, G. R. Goren-Muginstein and M. Bamberger, Department of Material Eng., Technion IIT, Haifa 32000, Israel .......................................................................... 134 Investigation of Die Cast Mg-Al Alloys with Additions of Sn, Si and Sb Dr. Per Bakke, Dr. Ketil Pettersen, Dr. Håkon Westengen, Hydro Aluminium a.s., Magnesium Competence Centre, Porsgrunn, Norway..................................................... 140 An Electron Microscopy Investigation on As-cast AZ91D Alloy Modified with Nitrogen A.A. Kaya, TUBITAK-MRC, MCTRI, Kocaeli, Turkey; O. Yucel, Istanbul Technical University, Istanbul, Turkey; D. Eliezer, Ben-Gurion University of The Negev Department of Materials Engineering Beer-Sheva Israel; E. Aghion, MRI, Beer-Sheva, Israel................................................................................. 151 Recrystallization Behaviour of Four Magnesium Alloys L.W.F. Mackenzie, F.J. Humphreys, G.W. Lorimer, Manchester Materials Science Centre, University of Manchester and UMIST, UK; K. Savage, T. Wilks, Magnesium Elektron Ltd, Manchester, UK .................................... 158 Dynamic Recrystallization of Mg-3Al-1Zn T. Laser, Ch. Hartig, R. Bormann, Technische Universität Hamburg-Harburg, 21073 Hamburg, Germany; J. Bohlen, D. Letzig, GKSS Forschungszentrum Geesthacht, 21502 Geesthacht, Germany ........................................................................................................................... 164 Mechanical Properties and Formability of PM Mg-Al Based Alloys K. Matsuzaki, K. Hatsukano, K.Hanada, M.Takahashi and T. Shimizu, National Institute of Advanced Industrial Science and Technology(AIST), AIST-EAST, 1-2 Namiki, TSUKUBA 305-3564, Japan ........................................................................ 170 Corrosion Properties of Conversion Plasma Coated Magnesium Alloys H. M. Nykyforchyn, M. D. Klapkiv, Karpenko Physico-Mechanical Institute of the National Academy of Sciences of Ukraine, Lviv, Ukraine; W. Dietzel, C. Blawert, GKSS-Forschungszentrum Geesthacht GmbH, Geesthacht, Germany ........................................................................................................................... 176 The effect of Variations in Melt Temperature on the Grain Structures in an AM60 Die Casting Hans I. Laukli, Otto Lohne, Lars Arnberg, Norwegian University of Science and Technology, Trondheim, Norway; Haavard Gjestland, Stian Sannes, Norsk Hydro ASA, Porsgrunn, Norway.................... 182

X Microstructural Characterisation of AZ91 Magnesium Alloy Andrzej Kiełbus, Maria Sozan´ska, Silesian University of Technology, Katowice, Poland; Lubomirc Ciz˙ek, VSB-TU Ostrava, Czech Republic ........................................................ 190 Microstructural Changes of AZ91 Magnesium Alloy after Heat Treatment Andrzej Kiełbus, Silesian University of Technology, Katowice, Poland; Lubomir Ciz˙ek, Libor Pawlica, VSB-TU Ostrava, Czech Republic................................. 196 Microstructure of Ultra Fine Grained Mg and Mg-10 wt.% Gd Prepared by HighPressure Torsion J. Cizek, I. Prochazka, I. Stulikova, B. Smola, R. Kuzel, V. Cherkaska, Charles University, Faculty of Mathematics and Physics, Prague, Czech Republic; I. Stulikova, B. Smola, Zentrum für Funktionswerkstoffe GmbH, Clausthal-Zellerfeld, Germany; R.K. Islamgaliev, O. Kulyasova, Institute of Physics of Advanced Materials, Ufa State Aviation Technical University, Ufa, Russia ............................................................ 202 Development of New Grain Structure and Tensile Properties Improving in a Hot Pressed and ECAP Processed ZK60 Magnesium Alloy A. Galiyev, R. Kaibyshev, M. Almakaev, Institute for Metals Superplasticity Problems, Ufa, Russia ....................................................................................................................... 208 Texture Formation and Texture Modelling of AZ31 Magnesium Wrought Alloy A. Styczynski, Ch. Hartig, R. Bormann, Arbeitsbereich Werkstoffphysik und – Technologie, Technische Universität Hamburg-Harburg, 21071 Hamburg, Germany; F. Kaiser, J. Bohlen, D. Letzig, GKSS-Forschungszentrum Geesthacht GmbH, 21502 Geesthacht, Germany............................................................................................ 214 Texture Development of AM20 Tensile Samples Under Load S.B. Yi, H.-G. Brokmeier, Institute of Materials Engineering, Technical University Clausthal, Clausthal-Zellerfeld, Germany; K.U. Kainer, T. Lippmann, Institute of Materials Engineering, GKSS-Research Center, Geesthacht, Germany ......................................... 220 Strengthening Phases in Extruded Quasicrystal Containing Mg-Zn-Y Alloys Alok Singh, A.P. Tsai, M. Nakamura, National Institute for Materials Science, Tsukuba 305-0047, Japan; M. Watanabe, A. KatoToyota, Motor Corporation, Toyota 471-8572, Japan................ 226 Grain Structure Characterization of AM60 Die Castings by Electron Backscatter Diffraction (EBSD) Measurements in SEM Hans I. Laukli, Otto Lohne, Lars Arnberg, Norwegian University of Science and Technology, Trondheim, Norway..................................................................................... 232 Possibilities of ECAP of Magnesium Alloy Katarzyna N. Braszczyn´ska, Institute of Materials Engineering, Technical University of Cze˛stochowa, Al. Armii Krajowej 19, 42-200 Cze˛stochowa, Poland .......................... 236

XI Electron Probe Micro Analysis of Sedimented Zirconium Particles in Magnesium Bruce Davis, Materials Science Department, Oxford University, now Magnesium Elektron, Manchester, UK; Keyna O’Reilly, Materials Science Department, Oxford University, John King, Magnesium Elektron, Manchester, UK............................................................................ 242 Superplasticity in Hot-rolled Magnesium Alloy AZ31 Sheets P. Zhang, J. Wendt, V. Šupik, J. Zhang, F. Holländer, Lehrstuhl Metallkunde und Werkstofftechnik, BTU-Cottbus, Cottbus, Germany ........................................................ 248 Extrusion of Magnesium Profiles and Microstructure Characterization Sören Müller, Ingwer Denks, Klaus Müller, Technical University Berlin, Berlin, Germany; Anke Pyzalla, Vienna Technical University, Vienna, Austria.......................................... 254 Investigating the Plastic Deformation Behaviour of Magnesium Alloy AZ31 by Plane Strain Forging of U and H Sections M.S.Yong, S.C.V.Lim, Singapore Institute of Manufacturing Technology, Singapore ......................................................................................................................... 260 Application of Isothermal Rolling for Improvement of Sheet Formability of a Magnesium Alloy A. Galiyev, R. Kaibyshev, D. Voronin, Institute for Metals Superplasticity Problems, Ufa, Russia ....................................................................................................................... 266 Influence of Al Content (<3%) on the Extrusion Limit Diagram Matthew R., BarnettDeakin University, Geelong, VIC, Australia; Ji-Yong Yao, University of Queensland, Brisbane, QLD, Australia; Chris Davies, Monash University, Clayton, VIC, Australia ............................................ 272 Comparison of Microstructure and Mechanical Properties of Indirect and Hydrostatic Extruded Magnesium Alloys J. Swiostek, J. Bohlen, D. Letzig, K.U. Kainer, GKSS Research Center, Institute for Materials Research, D-21502 Geesthacht, Germany ...................................................... 278 Improvements to the Rolling-Process in Order to Increase the Formability of Magnesium Wrought Alloys Friedrich-Wilhelm Bach, Michael Rodman, Mirko Schaper, Arne Rossberg, Institute of Materials Science (IW), University of Hanover, Germany; Eckart Doege, Gerrit Kurz, Institute for Metal Forming and Metal Forming Machine Tools (IFUM), University of Hanover, Germany.............................................. 285 The Effect of Strain-Hardening and Process Parameters on the Forgeability of Wrought Magnesium Alloys Johnson C.F. Chan, M.S. Yong, Singapore Institute of Manufacturing Technology; Johnson C.F. Chan, C.J. Tay, National University of Singapore.................................... 294

XII Superplasticity in Magnesium Alloys Processed by Different-Roll-Speeds Rolling Shangli Dong, Tsunemichi Imai, Ichinori Shigemastu, Naobumi Saito, Institute for Structural and Engineering Materials, National Institute of Advanced Industrial Science & Technology (AIST), Nagoya, 463-8560, Japan; Shangli Dong, Harbin Institute of Technology, Harbin, 150001, China......................... 300 Results of Tempered Hydroforming of Magnesium Sheets with Hydroforming Fluids Reimund Neugebauer, Michael Seifert, Fraunhofer Institute of Machine Tools and Forming Technology in Chemnitz, IWU, Germany ......................................................... 306 Bending Magnesium Extrusions R. Gradinger, M. Brandecker, H. Kilian, A. Wahlen, ARC Leichtmetallkompetenzzentrum Ranshofen GmbH (LKR), A-5282 Ranshofen, Austria....................................... 313 Drawing of Magnesium Wires at Ambient Temperature Jörg Eickemeyer, Albert Güth, Martina Falter, Ralph Opitz, Leibniz-Institut für Festkörper- und Werkstoffforschung Dresden, Institut für Metallische Werkstoffe, Dresden, Germany ........................................................................................................... 318 The MAGNEXTRUSCO Project: Process and Alloy Development for Hydrostatic Extrusion of Magnesium A.J. den Bakker, W.H. Sillekens, TNO Industrial Technology, Eindhoven, The Netherlands; J. Bohlen, K.U. Kainer, GKSS Research Center, Geesthacht, Germany; G. Barton, RWTH Aachen – Institute of Metal Forming (IBF), Aachen, Germany ........ 324 Microstructure and Mechanical Properties of Squeeze Cast AM60 Mg Composites Reinforced with AISI 304 Wires Jae Joong Kim, Jong Dae Lim, Hyundai Motors Company, Whasung, Korea; Min Soo Yoo, Nack J. Kim, Pohang University of Science and Technology (POSTECH), Pohang, Korea ........................................................................................... 331 Thermal Strain in Mg Composites Alexandra Rudajevová, Pavel Lukácˇ, Department of Metal Physics, Charles University, Prague, Czech Republic ................................................................................ 337 Synthesis of TiC Dispersed Composite in Mg-alloy Matrix D.W. Lee, Korea Institute of Machinery and Materials. Changwonn, Kyungnam, Republika Korea S.V. Aleksandrovkii, V.M. Sizyakov, D.V. Kutsenko, St. Petersburg Mining Institute (Technical University), Russia ......................................................................................... 342 Influence of Fibre Preform Permeability on Infiltration of Magnesium-Zinc Base Alloys M.S. Yong, R.I. Temple, Singapore Institute of Manufacturing Technology, Singapore; A.J. Clegg, Wolfson School of Mechanical and Manufacturing Engineering, Loughborough University, UK ................................................................................................... 348

XIII Magnesium Integral Foams – Production and Properties Markus Hirschmann, Matthias Lamm, Carolin Körner, Robert F. Singer, Department of Materials Science, Lehrstuhl Werkstoffkunde und Technologie der Metalle (WTM), Friedrich-Alexander-Universität Erlangen-Nürnberg, Erlangen, Germany ........................................................................................................................... 354 Creep of Reinforced and Unreinforced QE22 Magnesium Alloy M. Pahutova, V. Sklenicka, K. Kucharova, M. Svoboda, Institute of Physics of Materials, Academy of Sciences of the Czech Republic, CZ-616 62 Brno, Czech Republic; K.U. Kainer, Institute for Materials Research, Center for Magnesium Technology, GKSS Research Center, D-21502 Geesthacht, Germany ................................................ 360 Thermal Exposure Behaviour of Aluminum Borate Whisker Reinforced Magnesium Matrix Composites M. Y. Zheng, K. Wu, M. Liang, School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, P. R. China; S. Kamado, Y. Kojima, Department of Mechanical Engineering, Nagaoka University of Technology, Nagaoka 940-2188, Japan ...................................................................... 366 Creep Properties of Mg Dispersion-strengthened by Graphite Hans Ferkel, Institut für Werkstoffkunde und Werkstofftechnik, Technische Universität Clausthal, Agricolastr. 6, 38678 Clausthal-Zellerfeld, Germany................. 372 High Strength Si-O-C Ceramic Foams for the Reinforcement of Mg-Alloys J. Zeschky, T. Höfner, M. Scheffler, P. Greil, University of Erlangen-Nuremberg, Erlangen, Germany; J. Lo, Canada Centre for Mineral and Energy Technology, Ottawa, Ontario, Canada ............................................................................................................................. 378 Creep Behavior of Mg-Alloys AE42 and AZ91 Reinforced with Carbon Short Fibers S. Ataya, E. El-Magd, Department of Materials Science – LFW, RWTH, Aachen, Germany; S. Mielke, Kolbenschmidt, Neckarsulm, Germany........................................................... 384 High Performance Magnesium Composite Alloys with Mg2Si Dispersoids via Solid-state Synthesis Process Katsuyoshi Kondoh, Research Center for Advanced Science and Technology (RCAST), The University of Tokyo, Tokyo, Japan; Tatsuhiko Aizawa, Center for Collaborative Research (CCR), The University of Tokyo, Tokyo, Japan .................................................................................................................... 390 Fracture Toughness Behaviour of As-Cast and Long-term Annealed Pressure Die Cast Magnesium Alloys Doris Regener, Gabriele Dietze, Institute of Materials Engineering and Materials Testing, Otto-von Guericke-University Magdeburg, Germany ....................................... 396

XIV Influence of Strain Rate and Temperature on Deformation and Fracture Behaviour of Magnesium Alloy MgAl8Zn: Tests and Numerical Simulations E. El-Magd and M. Abouridouane,Department of Materials Science, RWTH Aachen Germany ........................................................................................................................... 402 Low-Temperature Superplasticity in Microcrystalline ZK60 and AZ91 Magnesium Alloys Processed by ECAP V. N. Chuvil’deev, M. Yu. Gryaznov, A. N. Sysoev, Research Physical & Technical Institute of Nizhny Novgorod State University, Nizhny Novgorod, Russia; T. G. Nieh, Lawrence Livermore National Laboratory, University of California, Livermore, USA; V. I. Kopylov, Physical - Technical Institute of National Academy of Science, Minsk, Belarus ............................................................................................................................. 409 Structural Durability and Fatigue Life Assessment for Magnesium Diecast Specimen Karl Ludwig Kotte, Sven Werdin, Dresden University of Technology, Institute of Solid State Mechanics, Dresden, Germany...................................................................... 415 Microstructure/Property Relationships of Three Mg-RE-Zn-Zr Alloys L. Gill, G.W. Lorimer, Manchester Materials Science Centre, Manchester, UK; P. Lyon, Magnesium Elektron Ltd, Manchester (UK) ..................................................... 421 Influence of Extrusion Methods on Microstructure and Mechanical Properties of AZ31 Magnesium Alloy J. Dzwonczyk, N. Hort, K. U. Kainer, GKSS Forschungszentrum GmbH, Geesthacht, Germany ........................................................................................................................... 427 Anisotropic Mechanical Properties of Extruded Magnesium Alloy AZ31 C.H.J. Davies, F. Xiong, and M. Battaini, Cooperative Research Centre for Cast Metals Manufacturing (CAST), School of Physics and Materials Engineering, Monash University, Victoria, Australia, 3800 ................................................................. 433 Mechanical Properties of Hot-rolled Magnesium AZ31 Sheets J. Wendt, V. Šupik, P. Zhang and J. Zhang, Chair of Physical Metallurgy and Materials Technology, Technical University of Brandenburg at Cottbus, Cottbus, Germany ........................................................................................................................... 439 High Cycle Fatigue of Magnesium Alloys Herwig Mayer, Bernhard Zettl, Universität für Bodenkultur, Vienna, Austria; Anton Stich, Hans-Günther Haldenwanger, Audi AG, Ingolstadt, Germany .................. 444 Fatigue Life Prediction of Magnesium Alloys for Structural Applications Yukio Miyashita, Zainuddin bin Sajuri, Takashi Umehara and Yoshiharu Mutoh, Nagaoka University of Technology, Nagaoka, Japan...................................................... 450 Correlation of Microstructure and Mechanical Properties of Rolled Magnesium Sheet AZ31 F. Kaiser, J. Bohlen, D. Letzig, K.U. Kainer, GKSS Research Center GmbH, Geesthacht, Germany; A. Styczynski, Ch. Hartig, Technical University Hamburg-Harburg, Hamburg, Germany ........................................................................................................................... 456

XV Effect of Warm Forming on Mechanical Properties of an AZ31 Magnesium Alloy Sheet H. Takuda, H. Fujimoto, D. Inoue, Kyoto University, Kyoto, Japan............................... 463 Stress State Dependent Deformability of Polycrystalline Magnesium T. Törköly, A. Eleöd, Gy. Ziaja, Budapest University of Technology and Economics, Budapest, Hungeria.......................................................................................................... 469 A Method to Determine the Damping Behaviour of Coatings on Cast Magnesium Alloy AZ91 J. Göken, D. Letzig, K.U. Kainer, GKSS Research Center, Institute for Materials Research, D-21502 Geesthacht, Germany....................................................................... 475 Microstructure and Mechanical Properties of Heat Treated Mg-Zn-Al Based Alloys Valérie Lemaire, Dominique Dubé, and Réal Tremblay, Université Laval, Québec, Canada ............................................................................................................................. 483 Analysis of Deformation and Fracture Mechanisms in Mg-AZ31 Sheet Material Gaofeng Quan , Jürgen Heerens, Institute for Materials Research, GKSS Research Centre, 21502 Geesthacht, Germany ............................................................................... 489 Unstable Plastic Deformation in Mg Alloys-post Relaxation Effect Zuzanka Trojanová, Pavel Lukácˇ, Faculty Mathematics and Physics, Charles University, Ke Karlovu 5, CZ-121 16 Praha 2, Czech Republic; Luboš Cˇížek, Technical University Ostrava, Trˇ. 17. listopadu 2172, CZ-708 00 Ostrava, Czech Republic .................................................................................................. 495 On the Reliability of Yield Strength Data Evaluated from as-cast Tensile Test Bars of Mg-alloys M. Wessén, H. Cao, Jönköping University, Jönköping, Sweden...................................... 501 Understanding the Corrosion Mechanism: a Framework for Improving the Performance of Magnesium Alloys Guangling Song and Andrej Atrens, CRC for Cast Metals Manufacturing (CAST), Division of Materials, The University of Queensland, Brisbane Qld Aus 4072, Australia ........................................................................................................................... 507 Potentiodynamic Studies of Some AZ and ZA Magnesium Alloys in Different Corrosive Media Mustafa Özgür Öteyaka, Anne-Marie Lafront, Edward Ghali and Réal Tremblay, Université Laval, Québec, Canada .................................................................................. 517 Plating of Magnesium – New Developments Peter Gregg, Franz Oberflächentechnik GmbH & Co. KG, Geretsried, Germany......... 524 Surface Treatment of Magnesium Substrates Michael Walter, AHC Oberflächentechnik Holding GmbH, Kerpen, Germany.............. 529

XVI Approach to Control the Corrosion of Magnesium by Alloying V. Kaese1, Volkswagen Group Research, Wolfsburg, formerly Institute for Materials Science and Biomedical Engineering Centre, University of Hanover; P.-T. Tai, Fr.-W. Bach, Institute for Materials Science, University of Hanover; H. Haferkamp, Institute for Materials Science, Biomedical Engineering Centre, University of Hanover; F. Witte, H. Windhagen, Hanover Medical School, Hanover, Germany......................... 534 Magnesium Alloys Protection by Microplasmic Processing Belozerov Valeriy, Isakov Sergey, Makhatilova Anna, Subbotina Valeriya, National Technical University “Kharkov Polytechnic Institute”, Kharkov; Isakova Tatiana – Innovation Center “MAGIC Solutions” ltd., Kharkov, Russia........................................ 540 Class A Surface Quality for Magnesium Diecastings Using Ceramic Precursor-based Coatings R. Gadow, D. Scherer, C. C. Stahr, University of Stuttgart, Institute for Manufacturing Technologies of Ceramic Components and Composites – IFKB, Stuttgart, Germany .... 544 Wear and Corrosion Resistant Ceramic-based Multilayer Coatings with Dry Lubrication Ability R. Gadow, D. Scherer, C. C. Stahr, University of Stuttgart, Institute for Manufacturing Technologies of Ceramic Components and Composites – IFKB, Stuttgart, Germany ........................................................................................................................... 548 New Generation of Keronite Process for Corrosion Protection of Magnesium S. Hutchins, P. Shashkov, V. Samsonov and A. Shatrov, Keronite Ltd, Granta Park, Great Abington, Cambridge CB1 6GP, UK..................................................................... 553 Corrosion Properties of PVD-coated Magnesium Die Cast Alloy AZ91D Investigation of the Corrosion Properties and the Corrosion Mechanisms of PVD Coated Magnesium Die Cast Alloy AZ91 for Tribological Applications H. Hoche, E. Broszeit, C. Berger, Staatliche Materialprüfungsanstalt und Fachgebiet und Institut für Werkstoffkunde an der Technischen Universität Darmstadt, Darmstadt; C. Blawert, GKSS Forschungszentrum Geesthacht GmbH, Geesthacht, Germany ........ 559 Application of Magnesium Alloys Coated with High Purity Mg for Engineering Usage A. Yamamoto, H. Tsubakino, Graduate School of Engineering, Himeji Institute of Technology, Himeji, Japan; R. Tamai, Graduate Student of Himeji Institute of Technology, Himeji, Japan .............. 568 Corrosion and Fatigue Behaviour of the Magnesium Die-Cast Alloy AZ91 hp after Surface Treatment C. Dindorf, C. Müller, Technische Universität Darmstadt, Germany............................. 574 Corrosion Behaviour in Artificially Corrosion-Oxidation treated Mg Alloys A. Yamamoto, H. Tsubakino, Graduate School of Engineering, Himeji Institute of Technology, Himeji, Japan; H. Inaoka, Graduate Student of Himeji Institute of Technology, Himeji, Japan............. 580

XVII Corrosion and Oxidation of Thixomolded Magnesium Alloys Claudia Pieper, Uwe Köster, Helena Alves, University of Dortmund, Germany; Isao Nakatsugawa, The Japan Steel Works, Ltd., Hiroshima, Japan .............................. 586 Plasma Electrochemical Synthesis of Protective Coatings on WE54 Magnesium Alloy A. Berkani, P. Skeldon, G.E. Thompson, Corrosion and Protection Centre, University of Manchester Institute of Science and Technology, P.O. Box 88, Manchester M60 1QD, UK; K. Shimizu, University Chemical Laboratory, Keio University, 4-1-1 Hiyoshi, Yokohama 223, Japan; H. Habazaki.. , Graduate Engineering School, Hokkaido University, N13 W8, Kita-ku, Sapporo 060-8628, Japan; H. Karimzadeh, Magnesium Elektron, P.O. Box 23, Swinton, Manchester M27 8DD, UK; K. Stevens, C. John, Poeton Industries Ltd., Eastern Avenue, Gloucester GL4 3DN, UK .................................................................................................................................... 592 Galvanic Corrosion of Magnesium Alloy AZ91 with Coated Elements E. Boese, J. Goellner, A. Heyn, Otto-von-Guericke Universität, Magdeburg, Germany; St. Eisenberg, Volkswagen AG, Wolfsburg, Germany ..................................................... 599 Corrosion Protection by Conductive Polymers for Magnesium Alloys A. Yfantis, I. Paloumpa, D. Schmeißer, BTU Cottbus, Department of Applied Physics and Sensors, Cottbus, Germany; D. Yfantis, NTU Athens, Department of Chemical Engineering, Athens, Greece ........... 605 Plasma Hydrogenation of Mg-based Alloy Films Under High-flux, Low Energy Ion Irradiation at Elevated Temperatures L. L. Pranevicius, Vytautas Magnus University, 8 Vileikos St., LT-3035 Kaunas, Litaric; L. L. Pranevicius, D. Milcius, Lithuanian Energy Institute, 3 Breslaujos St., LT-3035 Kaunas, Litaric; G. Thomas, Sandia National Laboratories, 537 Spirit Ridge Court, Reno, NV 89511, USA .................................................................................................................................. 611 The Corrosion Performance of Sealed KERONITE Coatings on AZ91D Magnesium Alloy S. Shrestha, A. Sturgeon, TWI Ltd, Granta Park, Great Abington, Cambridge CB1 6AL, UK; S. Hutchins, P. Shashkov, A. Shatrov, Keronite Ltd, Granta Park, Great Abington, Cambridge CB1 6GX, UK; .............................................................................................. 617 Corrosion Properties of PVD-coated Magnesium Die Cast Alloy AZ91D Contact Corrosion Properties of Differently PVD treated Magnesium Die Cast Alloy AZ91 H. Hoche, E. Broszeit, C. Berger, Staatliche Materialprüfungsanstalt und Fachgebiet und Institut für Werkstoffkunde an der Technischen Universität Darmstadt, Darmstadt, Germany; C.Blawert, GKSS Forschungszentrum Geesthacht GmbH, Geesthacht, Germany ......... 623

XVIII Corrosion Stress Relaxation in Magnesium Alloys Ya. Unigovski, L. Riber, A. Eliezer, E. M. Gutman, Ben-Gurion University of the Negev, Beer-Sheva, Israel ............................................................................................... 632 Effect of Small Ca-additions on the Corrosion Resistance of AS41-Mg Alloy Volkmar Neubert, Institut für Materialprüfung und Werkstofftechnik, Clausthal, Germany; Ashraf Bakkar, Dept. of Metallurgy & Materials Engg., Suez Canal University, Egypt................................................................................................................................. 638 Contribution to the Study of Electrochemical Corrosion of Alloy AZ91 in Selected Environment S. Lasek, L. Cˇ ížek, L. Pawlica, VŠB-TU Ostrava, Department of Materials Engineering, Ostrava, Czech Republic........................................................................................... 646 Study on Effects of RE, Ca Addition on Stress Corrosion Cracking Behaviour of Magnesium Alloys Z. L. Wei, Q .R. Chen, X. C. Guo, L. Yang, N. X Xiu, Y. W. Huang, Corrosion & Protection Center of Metals, Shanghai Institute of Microsystem and Information Technology, Chinese Academy of Sciences No.865 Changning Road, Shanghai, 200050, P.R. China ........................................................................................................ 649 Comparison of the Corrosion Behaviour Between Two Magnesium Alloys by Electrochemical Noise and Scanning Reference Electrode Technique Robert D. Klassen, Pierre R. Roberge, Department of Chemistry & Chemical Engineering, Royal Military College of Canada, Kingston, Canada; Anne-Marie Lafront, Mustafa Ö. Öteyaka, Edward Ghali, Department of Mining and Metallurgy, Laval University, Quebec, Canada .............................................................. 655 Role of Manganese in the Soldering Reaction in Magnesium High Pressure Die Casting C. Tang, M. Brandt, Industrial Research Institute Swinburne (IRIS), Swinburne University of Technology, Australia; M. Jahedi, CSIRO Manufacturing and Infrastructure Technology (CMIT), Australia ... 661 Extended Casting Simulation for Improved Magnesium Die Casting Erik Hepp, MAGMA Giessereitechnologie GmbH, Aachen, Germany; Otto Lohne, Norwegian University of Science and Technology (NTNU), Trondheim, Norway; Stian Sannes, Hydro Aluminum Research Center, Porsgrunn, Norway.......................... 669 Vertical Direct Chill (VDC) Casting of Magnesium – Optimized Casting Parameters and Safety Issues Franka Pravdic, Dietmar Leitlmeier, ARC Leichtmetallkompetenzzentrum Ranshofen GmbH (LKR), Austria; Christian Wögerer, Gerhard Traxler, ARC Seibersdorf research GmbH (ARCS), Austria; Herbert Hubbauer, HPI Ges.m.b.H, Austria ................................................................... 675

XIX Innovative Magnetohydrodynamical Effect Use in Casting Magnesium Alloys Viktor I. Dubodelov, Vladimir A. Seredenko, Viktor K. Pogorsky, Physico-Technological Institute of Metals & Alloys of NASU, 34/1 Vernadsky Ave, Kyiv-142, 03680, Ukraine; Vladimir G. Tkachenko, International Center for Electronic Materials Science and Applied Problems of Aerospace Technology, 3, Krzhyzhanovsky St, Kyiv-142, 03680, Ukraine................................................................................................................. 681 Magnesium New Rheocasting as an Industrial Process: An Update of Recent Developments Werner Fragner, Helmut Kaufmann, ARC Leichtmetallkompetenzzentrum Ranshofen GmbH (LKR), Ranshofen, Austria; Helmut Helmberger, Ralph Potzinger, TCG Unitech AG, Kirchdorf/Krems, Austria..... 688 Influence of Process Parameters on the Microstructure and Mechanical Properties of Magnesium Die Castings C. Pitsaris, T. Abbott, C.H.J. Davies, Monash University, Australia; G. Savage, CSIRO Division of Infrastructure and Manufacturing Technology, Australia ........................................................................................................................... 694 Residual Stresses in Thin-walled Hot Chamber Die Cast AZ91d; an Experimental Study Anders Eric Wollmar Jarfors, Singapore Institute of Manufacturing Technology, Singapore; Tan Ming Jen, Nanyang Technological University, Singapore; Jonas. Bäckman, University of Jönköping, Sweden......................................................... 700 A New Zirconium-rich Master Alloy for the Grain Refinement of Magnesium Alloys Ma Qian, D. H. StJohn, Division of Materials Engineering, University of Queensland, Brisbane, Qld 4072, Australia; M. T. Frost, Australian Magnesium Corporation Limited (AMC), PO Box 1364, Milton BC, Qld 4064, Australia ....................................................................................... 706 A Manufacturing Technology of AZ91-alloy Slurry for Semi Solid Forming Ji-Young Byun, Soon Il Kwon, Heon Phil Ha, and Jin-Kook Yoon, Div. of Materials, Korea Inst. of Sci. and Tech., #39-1 Haweolgok-Dong, Sungbuk-Ku, Seoul 136-791, Republic of Korea............................................................................................................. 713 Characterization and Removal of Die Lubricant Residue on Magnesium Die Castings Jan Ivar Skar, Robert Orr, Arne Henriksen, Norsk Hydro Research Park, Porsgrunn, Norway; Heike Thamerus, Jochen Caster, Klüber Lubrication, Munich, Germany ...................... 719 Magnesium Die Casting for High Performance Stian Sannes, Haavard Gjestland, Håkon Westengen, Magnesium Competence Centre, Hydro Research Park Porsgrunn, N-3907 Porsgrunn, Norway; Hans Ivar Laukli, Otto Lohne, Department of Materials Technology, NTNU, N-7491 Trondheim, Norway.......................................................................................................... 725

XX Influence of Direct Chill Casting Parameters on a Magnesium Alloy D. Ludlow, G.W. Lorimer, Manchester Materials Science Centre, Manchester, UK; J-C. Gebelin, W.D. Griffiths, University of Birmingham, Birmingham, UK; J. King, T. Wilks, K. Savage, Magnesium Elektron Ltd, Manchester, UK....................... 732 Injection Molding of Magnesium Alloys Andreas Lohmüller, Martin Scharrer, Ralf Jenning, Michael Hilbinger, Mark Hartmann, Robert F. Singer, Neue Materialen Fürth GmbH (NMF), Dr.-Mack Strasse 81, D-90762 Fürth, Germany................................................................................................. 738 Molten Metal Jets Lamination at Pouring into Casting Molds A. I. Kulinsky, JSC AVISMA Titanium-Magnesium Works, Berezniki, Russia ................ 744 Non-destructive Testing of Cast Magnesium Alloys by Computerized Tomography Dr. Bernhard Illerhaus, Dr. Jürgen Goebbels, Bundesanstalt für Materialforschung und -prüfung (BAM), Berlin, Germany; Dr. Norbert, Dr. Carsten Potzies, Hagen Frank, HortGKSS-Forschungszentrum GmbH, Geesthacht, Germany .......................................................................................... 746 Investment Casting of Magnesium M. Rosefort, S. Korte, A. Bührig-Polaczek, Giesserei-Institut RWTH Aachen, Germany ........................................................................................................................... 752 Cast Magnesium Alloys and Technological Processes for Casting Production I.Y. Mukhina, I.E. Kovalev, I.N. Fridlyander, I. N. Vasiliev, Federal State Unitary Enterprise “All-Russian Institute of Aviation Materials”, 17 Radio str., Moscow, 105005, Russia ................................................................................................................. 758 Rheo-Container-Process (RCP): New Semisolid Forming Method for Light Metal Alloys Julio Aguilar, Tilman Grimming, Andreas Bührig-Polaczek, Foundry Institute, Aachen University (RWTH), Aachen, Germany............................................................... 767 Thin Slab Casting as a New Possibility for Economic Magnesium Sheet Metal Production Heinz Palkowski, Lutz Wondraczek, Institut für Metallurgie, Technische Universität Clausthal, Clausthal-Zellerfeld, Germany....................................................................... 774 Duplex Process: Crucible Furnace – Salt-Heated Furnace is an Effective Process of Making Magnesium Alloys I.A. Barannik, I.M. Komelin, Titanium Institute, Zaporozhye, Ukraine .......................... 783 Development of a PAM-CRASH Material Model for Die Casting Alloys Lars Greve, Volkswagen AG, Wolfsburg, Germany ........................................................ 789 Integrated Model for Calculating Microstructural and Forming Parameters of Magnesium Alloy AZ31 during Hot Rolling Kawalla, R., Stolnikov, A., Technische Universität Bergakademie Freiberg, Germany Coung, N., Magnesium Flachprodukte GmbH, Freiberg, Germany ............................... 803

XXI Advanced Die Casting Simulation for a Systematical Development of Magnesium Structural Components Christoph Schendera, Simone Schwarz, Martin Möser, Fraunhofer IWM, Freiburg Germany; Thorsten Müller, Fraunhofer IFAM, Bremen, Germany.................................................. 811 Fatigue Simulation of Light Weight Components of Magnesium Ulrich May, Antje Berg-Pollack, Fraunhofer-Institut für Betriebsfestigkeit LBF, Darmstadt, Germany........................................................................................................ 817 Product Development Process Optimization by Linking Design, Numerical Simulation and Process Development Elke Lieven, Moritz Wuth, TAKATA-PETRI Engineering Aschaffenburg - Berlin - Elterlein - Ulm .............................................................................................................. 823 Calculation of Phase Dissolution and Homogenisation after Non-Equilibrium Solidification in MgAlZn Alloys Eva Cyrener, Eileen Clauss, Jörg Buchner, Westsächsische Hochschule Zwickau (FH), (University of Applied Sciences), Zwickau, Germany............................................ 835 Solidification Curves of AZ-Magnesium Alloys Determined by DSC Experiments and Heat-transfer Model (DSC-HTM) Djordje Mirkovic´, Joachim Gröbner, Rainer Schmid-Fetzer, Technical University of Clausthal, Institute of Metallurgy, Clausthal-Zellerfeld, Germany................................. 842 Numerical Simulation of Magnesium Alloy Sheet in Thermal Deep-drawing Process Shi-Hong Zhang, Kun Zhang, Zhong-Tang Wang, Chuan-Fu Yu, Yi Xu, Ke Yang; Institute of Metal Research, Chinese Academy of Sciences, Shenyang, P.R. China........ 848 Machinability Investigations of Fibre Reinforced Magnesium Klaus Weinert, Matthias Lange, Department of Machining Technology, University of Dortmund, Dortmund, Germany.................................................................................. 854 Thermal Spraying of Magnesium Alloys for Corrosion and Wear Protection Erich Lugscheider, Maria Parco, Materials Science Institute, Aachen University of technology, Aachen, Germany; K. U. Kainer, N. Hort, GKSS-Research Centre Geesthacht GmbH, Geesthacht, Germany ........................................................................................................................... 860 Welding of Magnesium Alloy AZ91C by CO2 Laser and by Friction Stir Welding P. Volovitch, M. Dhahri, J.-E. Masse, L. Barrallier, MécaSurf ENSAM, Aix en Provence; T. Baudin, Université Paris-Sud, Orsay, France; B. DaCosta, J.C. Goussain, Institut de Soudure, Yutz, France........................................ 869 Investigations of High-Productivity Methods for Joining Magnesium Sheet Jerry E. Gould and Wayne L. Chuko, Edison Welding Institute Columbus, OH, USA ... 875

XXII High Speed Minimum Quantity Lubrication Machining of Magnesium Franz Obermair (Sp), Günther Klammer, Profactor Produktionsforschungs GmbH, Steyr, Austria.................................................................................................................... 881 Safe and Economically Efficient Use of Coolants in Mechanical Processing of Magnesium Alloys Stefan Joksch, Oemeta Chemische Werke GmbH, Uetersen, Germany........................... 888 Safe Machining of Magnesium Parts by Cutting and Burnishing Operations B. Denkena, J. C. Becker, C. Podolsky, A. Kuhlmann, Institute of Production Engeneering and Machine Tools, University of Hannover, Schlosswender Str. 5, 30159 Hannover, Germany .............................................................................................. 895 Development of a Flux Free Soldering Technology for Magnesium Alloys Bernhard Wielage, Silke Mücklich, Institute of Composite Materials, University of Technology Chemnitz, Germany ...................................................................................... 902 Fatigue Strengh of Welded Wrought Magnesium Alloys B. Kleinpeter, M. Rethmeier, H. Wohlfahrt, K. Dilger, Institute for Joining and Welding, Langer Kamp 8, 38106 Braunschweig, Germany ............................................ 908 Friction Welding of Magnesium: Microstructural Characterization A.L. Bowles, N. Hort, A. Meyer, J. F. dos Santos and K. U. Kainer, GKSS Forschungszentrum GmbH, Geesthacht, Germany ............................................................................ 917 Module Design Concept for Lightweight Automotive Seats Young Jin Ko, Jae Joong Kim, Soon Chan Park, Jong Myung Kim, and Jong Dae Lim, Hyundai Motor Company, Whasung, Republic of Korea ................................................ 924 Cast Magnesium Safety Components - Energy Absorption Capacity Dørum, C., Hopperstad, O.S., Langseth, M., Structural Impact Laboratory (SIMLab) Department of Structural Engineering, Norwegian University of Science and Technology 7491 Trondheim, Norway; Lademo, O.-G., SINTEF Materials Technology, 7465 Trondheim, Norway; Sannes S., Norsk Hydro ASA, 3907 Porsgrunn, Norway ................................................. 930 Application of Mg Sheets in Car Body Structures F. Moll, M. Mekkaoui, S. Schumann, H. Friedrich, Volkswagen AG, Wolfsburg, Germany ........................................................................................................................... 936 Experimental Comparison of Automotive Oil Pans: Creep Resistant Magnesium Alloy vs. Aluminum Alloy P. Vert, Tesma International INC, Concord (Canada); X.P. Niu, Toral Cast-Light Metal Technologies, Concord (Canada); E. Aghion, Dead Sea Magnesium Ltd, Beer-Sheva (Israel); A. Stickler, Magna Advanced Technologies, Aurora (Canada)....................................... 943

XXIII Approaches to Use Magnesium as Structural Material in Car Body Volker Kaese, Lars Greve, Sven Jüttner, Martin Goede, Soenke Schumann, Horst Friedrich, Volkswagen Group Research, Wolfsburg, Germany; Werner Holl, Werner Ritter, AUDI Leichtbauzentrum, Neckarsulm, Germany .............. 949 Magnesium Sheets for Industrial Application S. Bräunig, M. Düring, H. Hartmann, B. Viehweger....................................................... 955 Magnesium in Automotive Lightweight Design Under Life Cycle Aspects Julia Pflieger, Matthias Fischer, Peter Eyerer, University of Stuttgart, IKP, Department for Life Cycle Engineering, Leinfelden-Echterdingen, Germany; Jürgen Stichling, PE Europe GmbH, Leinfelden-Echterdingen, Germany ..................... 962 Drawn Magnesium Wire for Robotic Welding and for Mechanical Applications Peter Baumgart (Sp), Drahtwerk Elisental W. Erdmann GmbH & Co., Neuenrade, Germany ........................................................................................................................... 968 Recycling Towards Quality Excellence: a Cost Efficient Solution Juan Gómez-Cordobés, Rauch Melting Technique, Gmunden, Austria .......................... 974 Application of Remelted Post Consumer Scrap for Structural Magnesium Parts Christiane Scharf, André Ditze, Technische Universität Clausthal, Institut für Metallurgie, 38678 Clausthal-Zellerfeld, Germany; Carsten Blawert, GKSS Forschungszentrum Geesthacht GmbH, Institut für Werkstoffforschung, 21502 Geesthacht, Germany ................................................................... 980 Techniques for Recycling of Magnesium Scrap G. Hanko, C. Lochbichler, W. Riederer, University of Leoben, Austria; G. Macher, ecka granules – non ferrum Metallpulver Gesellschaft; St. Georgen bei Salzburg; Austria.............................................................................................................. 988 Industrial Application of the AMCover Cover Gas System for Magnesium Melt Protection Simon Cashion, Rob Bailey, Australian Magnesium Corporation (AMC), Milton, BC, Australia; Nigel Ricketts, Cooperative Research Centre for Cast Metals Manufacturing (CAST), Brisbane, Australia ............................................................................................ 995 Environmental Friendly Protection System for Molten Magnesium F.-W. Bach, A. Karger, Ch. Pelz, Institute of Materials Science (IW), University of Hanover, Germany Linde AG, Unterschleißheim, Germany.................................... 1001 New Perspectives with Magnesium Sheet Peter Juchmann, Sébastien Wolff, Salzgitter Magnesium-Technologie GmbH, Salzgitter, Germany....................................................................................................... 1006 Experiment and Simulation for Optimisation of Production and Usage Qualification of Magnesium Components – Objectives and Results of the Fraunhofer Alliance »Magnesium Light Weight Construction« Johann Georg Blauel, Thomas Götz, Fraunhofer IWM, Freiburg, Germany; Jens Standfuss, Fraunhofer IWS, Dresden, Germany................................................... 1013

XXIV Solidification Morphology and Shrinkage Behavior of Mg-Alloys in Chill- and Sand Casting A. Hennings, E. Schaberger-Zimmermann, A. Bührig-Polaczek, Foundry Institute, RWTH Aachen, Germany .............................................................................................. 1020 InMak– Innovative Magnesium Combined Structures for Car Bodies J.-F. Lass, Institute of Materials Science, University of Hannover, Germany; A. Metz, M. Wappelhorst, Honsel GmbH & Co. KG., Meschede, Germany; E. Hombergsmeier, EADS, Ottobrunn, Germany; H. Lanzerath, Ford Forschungszentrum, Aachen, Germany; P. Baumgart, Drahtwerk Elisental W. Erdmann GmbH & Co, Neuenrade, Germany; P. Cordini, Laser Zentrum Hannover e.V., Hannover, Germany ................................. 1026 Research Activities on Magnesium Alloys at Seoul National University, Korea Kwang Seon Shin, School of Materials Science and Engineering, Research Institute of Advanced Materials, Seoul National University, Seoul 151-744, Korea ................. 1034 Mg-Chassis – Advanced Manufacturing Technology for Automotive Chassis Components through Extensible and Sustainable Use of Mg Alloys Alexandra Ziese, Gerald Rausch, Dirk Lehmhus, Thorsten Müller, Fraunhofer Institute for Manufacturing and Advanced Materials (IFAM), Bremen, Germany ...... 1040 The MAGNEXTRUSCO project: European Community Research on Hydrostatic Extrusion of Magnesium W.H. Sillekens, TNO Industrial Technology, Eindhoven, The Netherlands; J. Bohlen, GKSS Research Center, Geesthacht, Germany ........................................... 1046 Static Recrystallisation and Textural Changes in Warm Rolled Pure Magnesium Ravi Kumar Nadella, Indradev Samajdar and Günter Gottstein, Institut für Metallkunde und Metallphysik, RWTH Aachen, D-52056 Aachen, Germany ........ 1052 Author Index ................................................................................................................. 1059

1

Microstructural Design for Enhanced Elevated Temperature Properties in Sand-castable Magnesium Alloys C.J. Bettles and M.A. Gibson CRC for Cast Metals Manufacturing (CAST), CSIRO Manufacturing & Infrastructure Technology, Private Bag 33, Clayton South MDC, Clayton, VIC. 3169, Australia.

Abstract The low pressure and gravity casting processing techniques offer the component designer an opportunity to produce pieces with increased complexity over those able to be fabricated via the high pressure die casting route. The design of appropriate microstructures for the enhancement of elevated temperature properties is dependent on the chosen processing route, with the low pressure and gravity techniques allowing for post-cast manipulation of the microstructure through heat treatment. There are competing microstructural requirements for strength and creep resistance in casting alloys, and the optimised microstructure must, of necessity, be a compromise. The contributions from solid solution strengthening, grain boundaries and precipitation processes are described, with particular reference to elevated temperature magnesium alloys.

Introduction The relatively high cost of magnesium alloy feedstock dictates that high pressure die casting (HPDC) be the preferred processing route for automotive components. However, HPDC has its own inherent problems in that the castings are, from necessity, less complicated and contain a relatively high level of porosity when compared to sand cast counterparts. This latter effect restricts the choice of post-cast heat treatments and requires that the alloys develop their properties during the casting process and, further, the final properties of each casting will be dependent on the inherited defect structure. Gravity and low pressure casting processes may have higher production costs, but they have the advantages over HPDC of producing components of increased complexity, having better control over casting defects, and providing greater opportunities for alloy design through the tailoring of microstructure and post-cast heat treatment. The development strategy for gravity casting alloys is therefore very different to that for HPDC alloys. There are many different ways to influence the microstructural development of a particular alloy and the optimisation of a desirable microstructure requires consideration of all these aspects, tempered by the requirement to balance the mechanical properties in accordance with the specifications dictated by the final application. Figure 1 is a schematic representation of the alloy development process, showing the various aspects to be considered and their possible interactions. A separate, but often overriding, constraint to any

2 alloy development program is that the new alloy be commercially competitive in the market place. Precipitation hardening is the cornerstone of alloy development for the enhancement of mechanical properties in gravity cast components. An alloy is defined as precipitation or age-hardenable when its hardness or yield strength can be increased as a function of time under isothermal conditions through the solid state nucleation and growth of precipitating phases. In terms of a simple binary phase diagram (Figure 2(a)), an age-hardenable alloy is characterised by the presence of a high degree of solid solubility of a solute at an elevated temperature, which decreases sharply with decreasing temperature to some much lower level at room temperature. The most suitable systems are those in which this large solubility, as close to Cm as practicable, can be retained in super saturated solid solution (ssss) at room temperature ) with only moderate quench rates from the solution treatment temperature (Figure 2(b)). Process Control

Grain Boundaries

(Temperature/Time Cycle of Solution Treatment and Artificial Ageing)

(Control of Grain Size and Grain Boundary structure)

Promising Alloy System (Phase Equilibria, Solid-solution Strengthening, Multi-phase)

HEAT RESISTANT ALLOY

OPTIMISATION OF MICROSTRUCTURE

Structural Stability

Precipitate Structure (Shape, Size, Orientation and Distribution)

Trace Element Additions

Interface Control

Figure 1. Schematic representation of the alloy development process 700

500 Maximum Solid Solubility, Cm

500

400

400

Mg17Al12

300

a)

All Stable α

Moderate Quench

300 Saturated α + Mg17Al12 200 99% Transformed All Supersaturated α

100

200

100 0 Mg

Solution Treatment Temperature

α Temperature (oC)

Temperature (oC)

600

10

20

30

40

50

Weight Percent Aluminium

0 -3 10

60

b)

10-2

10-1

1

1% Transformed

10 1

10 2

10 3

10 4

Time (h)

Figure 2. (a) The magnesium-rich end of the binary Mg-Al phase diagram, showing the characteristics of an age-hardenable alloy system (after [1]) and (b) a schematic diagram showing typical Time-Temperature Transformation curves which may be used to predict appropriate heat treatment conditions.

3 For acceptable elevated temperature properties, it is also essential for the precipitating phases to have long-term stability at the operating temperature, and a good indication of thermal stability is the melting point of the precipitating phase (the higher the melting point, the greater the thermal stability). King [2] has listed the most likely age-hardenable magnesium systems, on the basis of the criteria listed above. Some twenty-eight binaries have been identified as candidate systems, but almost half of these would be considered commercially unattractive (i.e. the alloying element is extremely expensive and the automotive industry, in particular, would not accept the high cost penalty involved). Of the remainder, most, as straight binary systems, do not possess all the desirable attributes for an elevated temperature alloy, and microstructural manipulations of the type depicted in Figure 1 are therefore required in order to achieve the necessary property improvements. The addition of ternary and quarternary alloying elements increases the complexity of the phase interactions and leads to a many-fold increase in the number of possible alloying element combinations and heat treatment variations that can be employed to optimise the resulting microstructure. It is possible to predict the behaviour of these multi-component systems, using computational thermochemistry (with the Calphad method for instance [3]). The accuracy of such predictions is dependent on the quality of the available thermodynamic data, and the acquisition of this data is a long-term process. For magnesiumbased systems the available database is still relatively small, but Schmid-Fetzer and Gröbner [4] have clearly demonstrated the value of these computational techniques in identifying the most promising compositional envelopes in several alloy systems. However, until the database is much more comprehensive, most alloy development will continue to rely largely on an empirical approach, and it is the intention of this paper to highlight aspects of the individual contributions of the various microstructural components and heat treatments to the overall success of alloy design, with particular reference to magnesium alloys for elevated temperature applications.

Matrix Effects During heat treatment (natural or artificial ageing) after solutionising, the supersaturated solid solution decomposes to form a precipitating phase and a matrix phase, which retains an equilibrium concentration of homogeneously distributed solute atoms. The properties of both of these phases contribute to the mechanical behaviour of the alloy. For the matrix the solute atoms remaining in solid solution provide the capability of solute strengthening as defined by the relationship [5] of the form σy ∝ εs3/2 C1/2,

(1)

where εs is a misfit strain representing the size difference between the solute and solvent atoms and C is the solute concentration, which indicates that these misfit strains can be significant (generally by impeding dislocation motion). Limited information is available on the solid solution strengthening of magnesium, but it is known that the most efficient strengtheners are Zn, Al, Cd, Tl and In [6]. Suzuki et al. [7] have shown that additions of low levels of Zn (<0.1wt.%) to Mg-4wt.%Y can also induce a strengthening effect by decreasing the stacking fault energy of the magnesium matrix, which makes the slip of adislocations more difficult.

4 Solute atoms in solid solution, if distributed homogeneously, can also have an effect on the lattice parameters of the matrix magnesium. This effect is solute dependent, being a function of the solute radius and the electron valence. Busk [8] has determined the effect of several element additions on both the a and c/a values for magnesium. This effect can be particularly important where the primary precipitating phase has limited coherency with the matrix. An appropriate change to the lattice parameter of the matrix can lead to improved coherency and therefore an increased strengthening effect during age-hardening.

Precipitate Structure On a microscopic level, precipitates within a structure impede the motion of dislocations. The macroscopic manifestation of this effect is an improvement in the mechanical behaviour at both room and elevated temperatures. The precipitate-dislocation interactions can be divided into two broad groups, depending on the precipitate type: 1. Non-deformable particles in the matrix, and 2. Coherent precipitates, which can be sheared by dislocations in the matrix. The improvements in strength due to these two precipitate types can both be estimated from known theoretical relationships. For alloys with precipitates in the first group, the Orowan equation [9], or a variant of same, may be invoked: ⎛ Gb Δτ o = 2 ⎜ ⎝ 4π 1 − ν

⎞ ⎛ 1 ⎞ ⎛ dp ⎞ ⎟ ⎜ ⎟ ⎜ ln ⎟ ⎠ ⎝ λ ⎠ ⎝ ro ⎠

(2)

where the critical factors are λ, the interparticle spacing, and dp, the mean diameter of the precipitates (assumed to be spherical). A similar equation exists for the effect of shearable particles: ⎛ 2 ⎞⎛ 1 ⎞⎛ F ⎞ Δτ = ⎜ ⎟⎜ ⎟⎜ ⎟ ⎝ b Γ ⎠⎝ λ ⎠⎝ 2 ⎠

3/2

(3)

where Γ is the dislocation line tension and F is a measure of the precipitate resistance to shearing. It is clear that, in both cases, the strengthening effect may be influenced by the size of the particles and the interparticle spacing. The interparticle spacing is dependent upon the number of nucleation sites, and the shape and orientation of the precipitates with respect to the matrix. The first consideration therefore for precipitation hardening is the ease of nucleation of the precipitates, as a high number of nuclei will lead to a small interparticle spacing. Nucleation is initiated by the formation of small unstable embryo clusters, either at lattice defect points or as a consequence of similar atoms coming together to reduce the overall energy of the system. Some of these clusters will reach a critical size, at which time they become thermodynamically stable, and can continue to grow and possibly undergo solidstate transformations. It is generally during this growth phase that the precipitate crystal structures become independent of the matrix structure. Nucleation may be either homogeneous, in which case it occurs anywhere within the matrix and is temperature dependent, or heterogeneous, where a limited number of nucleation sites are located at some imperfection in the structure (impurity atoms or grain boundaries for instance). It is important that there be as many nucleation sites as possible, not only to ensure a high number densi-

5 ty of precipitates but also to aid the growth kinetics. Growth of the precipitates from these nuclei may be controlled by the condition of the interface (and the atom-exchange processes that may take place there) or by diffusion of the solute atoms through the parent lattice. Short diffusion distances will aid growth. The first phases to nucleate may be metastable in nature. These phases are usually coherent with the matrix, and can lead to exceptional age-hardening. GP zones are examples of such precipitation, being stable clusters of solute atoms, which have the same structure as the parent material. As ageing progresses, these zones may act as nucleation sites for either metastable or stable precipitating phases and, ultimately, stable equilibrium phases will form. These may still have some coherency with the matrix. GP zones are common first precipitating phases in aluminium alloys, and lead to a significant age-hardening response. However, GP zones are far less common in magnesium-based systems, being reported in some Mg-Rare Earth (RE) binary alloys and QE22. Some evidence of GP zone formation in Mg-Zn binary alloys has also been published [10]. Consequently, in magnesium alloys age-hardening is usually achieved through the growth of semi-coherent metastable precipitates, such as in the Mg-Y system: Mgssss → β" → β'→ β (Mg24Y5).

(4)

Coherency is extremely important, with similar lattice spacings existing in both the matrix and the precipitate and only a small amount of strain accommodation. This strain field in the matrix provides an additional obstacle to dislocation motion, effectively increasing the size of the obstacle. The semi-coherent precipitates, by definition, will have some interplanar spacings in common with those in the matrix. It is this matching which influences the preferred habit planes of the precipitate within the matrix. These habit planes will affect both the interparticle spacing and the ability of dislocations to climb around the obstacles. Nie and Muddle [11] have modified the Orowan equation for hexagonal materials and have shown that it is preferable, for improved strength properties, to have precipitates growing on prismatic and pyramidal planes. The shape of the particles can also have a significant effect on the interparticle spacing. In magnesium alloys, slip at low temperatures most easily occurs on the basal plane, and precipitates lying parallel to the basal plane will have limited effect on the impedance of dislocations. Nie and Muddle also examined different precipitate shapes in combination with various habit planes, and have predicted that the greatest strengthening will come from plate shaped precipitates on prismatic or pyramidal planes. The precipitates present at peak hardness may still be metastable or they may be equilibrium phases. Either way, for superior elevated temperature properties, it is essential that the microstructure generated is thermally stable with time. The system is dynamic, and the microstructure will change in an effort to reduce its overall energy. In terms of interfacial energies at precipitate-matrix boundaries, larger particles are favoured over smaller ones and as exposure at elevated temperature continues there will be preferential growth of the larger precipitates and dissolution of the finer precipitates. As this growth process continues, the volume fraction may remain the same but the number of precipitates will decrease and the interparticle spacing will increase. This coarsening process can be minimised by reducing the size difference between the largest and smallest particles. This is achieved through controlled nucleation such that the nucleation sites are present everywhere at the same time and that the preliminary clusters of solute atoms all reach critical size at the

6 same time. If the distribution of these clusters is homogeneous the growth rate of all precipitates will be the same and, at peak hardness, the size distribution will be very narrow and the driving force for preferential growth/dissolution will be very small. In addition, the presence of trace elements that segregate to the precipitate/matrix interface can insulate the particle against coarsening to enhance microstructural stability. The discussion thus far has considered the role of precipitates in strengthening an alloy. It does not necessarily follow, however, that higher yield strengths at elevated temperatures will lead to improved creep resistance. Nor is it necessarily true that the precipitate orientations and morphologies that provide the best strength improvements will result in the greatest creep resistance. In fact Smola et al. [12] have shown that in some Mg-MnSc-RE alloys which have plate-shaped precipitates on the basal plane, the creep resistance at temperatures up to 350 °C is excellent. The plates are acting as obstacles to both cross slipping basal dislocations and non-basal prismatic slip. Similar improvements to the creep performance of a Mg-3wt.% Nd-1.35wt.% Zn alloy over that of the binary Mg3wt.% Nd alloy have been associated with the formation of a large number of fine, stable precipitates lying on the basal plane in the former alloy that are absent in the latter [13].

Modifications to the Precipitates and the Precipitation Process On a macroscopic level, appropriate ternary additions (at trace or greater levels) may increase the peak hardness value, enhance the kinetics of age-hardening, or stabilise the precipitate microstructure to retard overageing. The mechanisms by which these alloying additions effect these changes are not well understood. At the trace addition level, it is known that these atoms can effectively act as nucleation sites, creating a strain field of opposite sign to that of the primary solute, attracting the primary solute atoms and thereby aiding the formation of precipitating phases. Trace additions may also change the precipitate morphology or habit planes during precipitation, by affecting the surface chemistry or interfacial energies. The modified interfacial effects may also affect the kinetics of overageing since this is influenced by free energy minimisation and diffusion processes. Such trace element effects have been clearly demonstrated in AZ91, where the addition of trace amounts of Cu increased the number of precipitates in the matrix by a factor of ten [14]. In another example, the typical "Chinese Script" morphology of Mg2Si is extremely detrimental to the mechanical properties of alloys containing this phase, but the addition of Ca or Sb induces the Mg2Si to form with a more refined blocky morphology, which leads to an improvement in the mechanical properties [15, 16]. Moreover, recent work by the authors [17], on the development of low cost creep resistant alloys based on the Mg-Zn system have revealed that the use of a proprietary trace addition can have a substantial positive influence on the precipitation process during artificial ageing. It has been determined that not only can the initial kinetics of the age hardening process be accelerated but also the onset of overageing can be significantly delayed, both of which are highly desirable in a low cost creep resistant alloy. It has also been found that the resulting microstructure of the modified alloy differs dramatically from that of the unmodified alloy, in that it possesses a more complex and much refined precipitate microstructure, with a greater number density of particles, compare Figures 3(b) and 3(d) with Figures 3(a) and 3(c).

7 It is not always possible to achieve sufficient improvements in mechanical behaviour and thermal stability by simply manipulating the precipitation mechanism of the binary precipitating phase. This is particularly the case with the Mg-Al(-Zn) system, where the presence of the Mg17Al12 phase is not conducive to improved elevated temperature behaviour. In these cases it is necessary to consider the precipitation of ternary phases, which generally requires the use of larger quantities of the alloying elements. Zhang et al. [18] increased the zinc content significantly, producing the ZA range of alloys. Mendis et al. [19] have characterised the ternary phases that form in this system, identifying an icosahedral phase Mg9Zn4Al3. They have also shown that by adding Ca to the alloy the icosahedral phase is replaced by a cubic phase, and this leads to greater thermal stability and improved creep resistance.

Grain Boundary and Grain Size Effects The grain refinement of casting alloys (either through control of cooling rate or the addition of grain refining agents) has long been recognised as a key factor in achieving desirable strength properties in a material. The Hall-Petch [20, 21,] relationship, which describes the functional relationship between the flow stress, σf, and the grain diameter, l, thus: σf = σo + k l -½,

(5)

has been applied almost universally to granular materials, and magnesium alloys are no exception. However, in precipitation-hardened alloys with a number of complex microstructural features, the yield strength – l -½ relationship may not be appropriate when l is taken to be the overall grain size. This has been found to be the case for AM-SC1 [22] where the microstructure contains intragranular precipitation, bulk boundary phases and precipitation along the boundaries and for which the yield strength increases linearly as the grain size decreases. Nonetheless, at its simplest level, the yield strength of a material can be assumed to increase as the grain size decreases. The elevated temperature creep behaviour of an alloy can also be affected by the overall grain size, and the conventional power law creep equation has been modified to include a grain size term: p

n

Q

⎛ b ⎞ ⎛σ ⎞ − ε& = A ⎜ ⎟ ⎜ ⎟ e RT ⎝d⎠ ⎝G⎠

(6)

where the values of p and n are related to the dominant creep mechanism [23]. Grain boundary sliding is an important creep mechanism, particularly in hexagonal materials, and in general, materials with a larger grain size show a lower steady state creep rate than those with a finer grain size. A similar effect occurs in materials where diffusional creep is the dominant mechanism – the rate of diffusional creep increases rapidly as the number of diffusion paths is increased (i.e. as the grain size decreases). Clearly, on the basis of these two grain size relationships (equations 5 & 6) the optimum grain size for a given material will depend entirely on which aspect of the mechanical behaviour is the most critical to the application under consideration and, of necessity, will generally be a compromise between the strength and creep properties.

8

a)

500 nm

b)

c

c

β1

a3

a2

a3

a2

β2 c)

a1

d)

a1

Figure 3. a) and b) Bright field electron micrographs of a) Mg-4wt.% Zn and b) Mg-4wt.% Zn modified with a 4 proprietary trace alloying addition (after ~10 ks at 177 °C), both taken with the electron beam parallel to the c-axis of the magnesium matrix. c) and d) schematic representations of the precipitate morphology and habit with respect to the Mg matrix for the unmodified (a) and the modified (b) alloys respectively.

The nature of the grain boundary itself can influence the creep behaviour of a material, and the detrimental effects of a refined grain structure may be offset by modifications to the grain boundary itself. Such modifications can be achieved through appropriate alloying additions. Cavity formation is generally the precursor to creep failure. Cavities form as a consequence of a stress concentration at some point, often associated with triple points, which are the sites of greatest disorder and inherent weakness. The presence of a stable, hard intermetallic phase at triple points, for instance, may inhibit the sliding of one grain over another, and reduce the opportunities for cavity formation at the triple points. The presence of particles in the grain boundaries may inhibit both the nucleation and growth of these cavities. Large, closely spaced particles, which have been formed during the initial solidification process (Figure 4(a)) may restrict grain boundary sliding but they can, especially if they are not coherent with the matrix, act as sites for stress concentration and aid the nucleation of voids and cavities. On the other hand, finely distributed, coherent precipitates within the boundaries (Figure 4(b)) are far more effective at inhibiting grain boundary sliding. These precipitates may form along the grain boundary dislocations, effectively pinning them in place. General fine precipitation in the boundary will interfere with any subsequent dislocation motion along the boundaries. The presence of precipitates in the boundaries may also slow diffusion in and around the boundary. For precipitation to occur within the grain boundaries during heat treatment, diffusion of solute atoms to the boundaries (and the equivalent migration of vacancies in the opposite

9

Grain 1

Grain 2

Grain 1 Grain 2

a)

500nm

b)

Figure 4. Grain boundaries in a T6 heat treated Mg-Rare Earth alloy showing (a) blocky precipitates formed during solidification that have not redissolved during solution heat treatment and (b) finely distributed coherent precipitates formed during the ageing process.

direction) must have occurred. Subsequent diffusion along the boundaries during service at elevated temperatures, which also requires the presence of vacancies, is therefore inhibited when pre-precipitation has taken place. Trace additions which locate at grain boundaries may also, directly or indirectly, affect the elevated temperature behaviour of an alloy. An example of this has been reported in AZ91 [24] where, in the T6 condition, the alloy contains precipitates of Mg17Al12 in two forms – platelets of continuous precipitation within the grains lying parallel to the basal plane, and lamellae of discontinuous precipitation originating at grain boundaries and sweeping into the grain interior. As nucleation of the platelets of Mg17Al12 is difficult, the first morphology to precipitate is the discontinuous lamellae from the grain boundaries. It is this form of the precipitate phase that is believed to be detrimental to the elevated temperature creep resistance of AZ91. The addition of trace amounts of gold altered the grain boundary environment and completely suppressed the formation of the discontinuous precipitation during the T6 heat treatment, resulting in a delay (by a factor of three) in the onset of tertiary creep in the Au-modified alloy over that of unmodified AZ91.

Process Control The precipitation reactions described thus far are all both time and temperature dependent. Optimising the heat treatment cycle requires consideration of both the solution treatment stage and the subsequent ageing stage (natural or artificial), together with the commercial constraints imposed by a production unit. The solution treatment temperature may, in the first instance, be selected from the equilibrium phase diagram. However, it is generally not sufficient just to choose a temperature that will lead to a single-phase material. The second stage of a heat treatment, where agehardening is achieved, may depend on the vacancy concentration in the super saturated solid solution (to reduce strain misfits, to aid in diffusion or to provide heterogeneous nucleation sites). In such cases the temperature should be as high as practicable, as the concentration of vacancies increases as the temperature increases. However, acting against this is the driving force for grain growth. Grain boundaries are sites of high energy and

10 therefore the natural tendency is to lower the total free energy by reducing the grain boundary area. As the temperature increases, this process becomes more energetically favourable. It is sometimes possible to suppress grain growth through trace additions, but the usual compromise solution is to optimise time and temperature to avoid substantial grain growth. The quenching rate from the solution treatment temperature must be sufficiently fast to both retain the solute in solid solution and maintain the highest vacancy concentration possible. A Time-Temperature-Transformation curve (see Figure 2(b)) gives a clear indication of whether or not this is possible and, additionally, will indicate the most appropriate temperatures and times for artificial ageing. Unfortunately, these are not readily available for most magnesium alloy systems, and ageing curves must be determined experimentally. In general terms the lower the ageing temperature the slower the hardening response but the finer the microstructure. The solution treatment can be a single or a multi-staged procedure, depending on the phases present. A microstructure containing both low and high melting point phases may need a dwell time at a temperature lower than the lowest melting point to take this phase back into solid solution before exposure at a higher temperature to achieve full solution treatment without incipient melting. Magnesium alloys are generally solution treated in an inert atmosphere to avoid oxidation but it is possible to use a reactive atmosphere during the heat treatment. For example, the alloy ZE63 is treated in a hydrogen atmosphere as the formation of hydrides within the matrix is an integral component of the strengthening process. The ageing schedule following solution treatment may also be a single or multi-stage process, depending on the metastable phases that precipitate out, and the nucleation mechanism for the most effective strengthening precipitate. Most of the aluminium and rare earth-containing alloys only require a single ageing treatment to achieve the desired age hardening response. Binary magnesium-zinc alloys, however, benefit from a double agehardening process [25]. Initial ageing is carried out at a temperature below 75 °C, to allow for the formation of GP zones, and a second stage then follows at 170 °C. The microstructure in these double-aged alloys is reported to contain a greater number of the primary strengthening phase, β1'. The precipitates are also much finer in size, than those achieved after a single stage ageing at the higher temperature.

Concluding Remarks An alloy microstructure is multi-faceted, and the development of the desired mechanical properties can require the manipulation of more than one of these microstructural features. The interactions between alloy chemistry, time and temperature are extremely complex and it is clear that there is no simple recipe for the development of alloys for elevated temperature applications. The approach to magnesium alloy development remains largely an empirical process, however, when due consideration is given to the various aspects detailed above substantial advancements in material properties can be achieved.

11

Acknowledgements This work was funded in part by the Australian Cooperative Research Centre for Cast Metals Manufacturing. The CRC for Cast Metals Manufacturing (CAST) was established under and is supported by the Australian Government’s Cooperative Research Centres Scheme.

References [1] Phase Diagrams of Binary Magnesium Alloys (Eds.: A. A. Nayeb-Hashemi and J. B. Clark), 1988, 18. [2] J. F. King in Magnesium Alloys and Their Applications (Eds.: B. L. Mordike and K. U. Kainer), 1998, 37-48. [3] L. Kaufman and H. Bernstein “Computer Calculation of Phase Diagrams with Special Reference to Refractory Materials”, Academic Press, New York, 1970. [4] R. Schmid-Fetzer and J. Gröbner, Advanced Engineering Materials 3(12) 2001, 947-961. [5] M. F. Ashby and D. R. H. Jones, Engineering Materials 2, Pergamon Press 1988, Chapter 10. [6] E. D. Levine, W. F. Sheely and R. R. Nash, Trans. AIME, 1959, 215, p 521 [7] M. Suzuki, T. Kimura, J. Koike and K. Maruyama, Mater. Sci. Forum, 426-432, 2003, 593-598. [8] R. S. Busk, Journal of Metals 188 (1950) 1460-1464. [9] E. Orowan in symposium on Internal Stresses in Metals and Alloys, 1948, p 451. [10] G. W. Lorimer in Magnesium Technology (Eds.: C. Baker, G. W. Lorimer and W. Unsworth), 1986, 47-53. [11] J. F. Nie and B. C. Muddle in Magnesium Alloys and Their Applications (Eds.: B. L. Mordike and K. U. Kainer), 1998, 229-234. [12] B. Smola, I. Stulikova, F. von Buch and B. L. Mordike, Mat. Sci. Eng. A324, 2002, 113-117. [13] R. S. Wilson, C. J. Bettles, B. C. Muddle and J. F. Nie in Proceedings of the Light Metals Technology Conference 2003 Ed. A. Dahle (in press). [14] C. J. Bettles in Magnesium Alloys and Their Applications (Eds.: B. L. Mordike and K. U. Kainer), 1998, 265-270. [15] M. O. Pekguleryuz and M. M. Avedesian in Magnesium Alloys and their Applications (Eds.: B. L. Mordike and F. Hehmann), 1992, 213-220. [16] G. Y. Yuan, Z. L. Liu, Q. D. Wang and W. J. Ding, Mater. Lett., 56, 2002, 53-58. [17] C. J. Bettles, M. A. Gibson and K. Venkatesan, to be published. [18] Z. Zhang, A. Couture and A. Luo in Magnesium Alloys and Their Applications (Eds.: B. L. Mordike and K. U. Kainer), 1998, 289-294. [19] C. Mendis, L. Bourgeois, B. C. Muddle and J. F. in Magnesium Technology 2003 (Ed.: H. Kaplan), TMS 2003, 183-188. [20] E. O. Hall, Proc. Phys. Soc. B64, 1951, 747-752. [21] N. J. Petch, J. Iron Steel Inst. 174, 1953, 25. [22] G. Dunlop, C. J. Bettles, J. R. Griffiths, K. Venkatesan and Li Zheng in Magnesium Alloys and Their Applications, 2003 (in press).

12

Thermodynamic Database for Mg-alloys: Progress and Application to Mg-Al-Ca-Si Alloy Development Rainer Schmid-Fetzer and Joachim Gröbner University of Clausthal, Institute of Metallurgy, Clausthal-Zellerfeld, Germany

1

Introduction

Rare earths (RE) additions can improve creep resistance of magnesium alloys by formation of precipitates. Some attempts were made in the last decades to replace the RE metals by calcium. Pekguleryuz and Baril [1] showed that diecast Mg-Al-Ca alloys offer creep resistance, tensile strength and corrosion resistance comparable to commercially used Mgalloys with RE additions, like AE42. These properties can be achieved with much lower costs. The stability of these Al2Ca precipitates is comparable to the RE precipitates. Von Buch et al. [2] described the development of new commercial magnesium alloys based on AZ91 with Calcium and minor RE (Ce mischmetal), Sr and Si additions. Calcium additions for improvement of the creep properties offer several additional advantages compared to cost intensive RE additions. It protects the melt surface combined with less slag formation and additive loss. The Al content can be high compared to Si containing alloys, which leads to better castability. The resulting so- called MRI alloys show lower secondary creep rates and higher tensile strengh than AZ91combined with comparable castability and good melt handling [2]. A better understanding of such new alloy developments and further modifications requires a sound knowlegde of their constitution as detailed below. As an example the predicted impact of Ca-addition to AS alloys is discussed and also the replacement of Ce in AE alloys by Ca.

2

Magnesium Database Development

Phase diagrams of multicomponent Mg-systems provide important information for alloying behavior. This information, plus the wealth of additional information accessible by thermodynamic calculations, forms a strong basis for the focused development of magnesium alloys [3]. The strength of this approach lies in the generation of consistent and quantitative – though easily perceived data – for example by “alloy tracking” through various solidification modes for any selected multi-component (Ⰷ 3) Mg-alloy: The kind of precipitating phases, their sequence and amounts are given as well as chemical activities for possible reactions with the environment or heat evolution data. This approach requires the use of modern thermochemical software [4] and – most importantly – a consistent thermodynamic database for multi-component Mg-alloys. This database is being constructed in an ongoing effort since many years in our group at TU Clausthal [3]. In order to use the Calphad approach as a tool in magnesium alloy design, a

13 high quality thermodynamic database for Mg-alloys is needed. Development of a reliable thermodynamic database for multicomponent alloys requires a combination of experiments and computational thermochemistry with data from alloy application. Since numerous binary and ternary subsystems have to be treated and validated with key experiments before multicomponent alloys can be calculated reliably, this development becomes a long-term project. Our thermodynamic database currently includes the elements Ag, Al, Ca, Ce, Cu, Gd, Li, Mg, Mn, Nd, Sc, Si, Sr, Y, Zn, Zr. The binary systems constitute the core of this 16-component system. Figure 1 gives an overview of the necessary systems and the current state of process. Given bold in the first column are the included elements. 120 independent binary systems are necessary as basis for a 16-component system. The thermodynamic descriptions of the most important binary system were already created mainly by our group and shaded gray in Figure 1. Also shown are some selected quaternary aim systems Mg-Al-Ca-Si, Mg-Al-Li-Si and Mg-Mn-Sc-RE (with RE = Ce, Gd, Sc and Y). It is important to note that based on reliable binary data and partly necessary ternary data a reasonable calculation of real multicomponent systems can be usually performed by extrapolation. Ag-Al Ag-Ca Ag-Ce Ag-Cu Ag-Gd Ag-Li Ag-Mg Ag-Mn Ag-Nd Ag-Sc Ag-Si Ag-Sr Ag-Y Ag-Zn Al-Ca Al-Ce Al-Cu Al-Gd Al-Li Al-Mg Al-Mn Al-Nd Al-Sc Al-Si Al-Sr Al-Y Al-Zn Al-Zr Ag-Zr Ca-Ce Ca-Cu Ca-Gd Ca-Li Ca-Mg Ca-Mn Ca-Nd Ca-Sc Ca-Si Ca-Sr Ca-Y Ca-Zn Ca-Zr Ce-Cu Ce-Gd Ce-Li Ce-Mg Ce-Mn Ce-Nd Ce-Sc Ce-Si Ce-Sr Ce-Y Ce-Zn Ce-Zr Cu-Gd Cu-Li Cu-Mg Cu-Mn Cu-Nd Cu-Sc Cu-Si Cu-Sr Cu-Y Cu-Zn Cu-Zr Si Gd-Li Gd-Mg Gd-Mn Gd-Nd Gd-Sc Gd-Si Gd-Sr Gd-Y Gd-Zn Gd-Zr Li-Mg Li-Mn Li-Nd Li-Sc Li-Si Li-Sr Li-Y Li-Zn Li-Zr Mg-Mn Mg-Nd Mg-Sc Mg-Si Mg-Sr Mg-Y Mg-Zn Mg-Zr Mn-Nd Mn-Sc Mn-Si Mn-Sr Mn-Y Mn-Zn Mn-Zr Nd-Sc Nd-Si Nd-Sr Nd-Y Nd-Zn Nd-Zr Mg Sc-Si Sc-Sr Sc-Y Sc-Zn Sc-Zr Si -Sr Si-Y Si-Zn Si-Zr Sr-Y Sr-Zn Sr-Zr Al Y-Zn Y-Zr RE Zn-Zr

Si

Li

Ca

Mg Mg

Al

multicomponent

Figure 1: The binary systems and selected ternary systems form the basis for multi-component applications of the Mg-database. Three quaternary systems are shown as examples.

In Mg-systems the experimental literature data were very sparse in both, phase diagram data and thermodynamic properties. Several experimental methods are being applied in our group to produce a sufficient experimental basis for a multicomponent Mg-database. For preparation of highly reactive and evaporating alloys containing Mg, Li and Ca the levitation melting technique was adapted. The main advantages are crucible free melting, suppressing of evaporation with protecting gas overpressure and visual control of energy supply to avoid overheating. Other preparation methods include arc and electron beam melting for extremely high melting intermetallic phases. During thermal analysis (DTA) the alloys were protected in a sealed tantalum crucible. For phase and crystal structure analysis X-ray diffractometry at room and high temperature, electron microscopy and metallography are also used in a combined approach.

14 Thermodynamic descriptions supported by experimental work of several subsystems were published recently for Al-Ca [5,6], Al-Li-Si [7,8], Ca-Li-Mg [9], Al-Ce-Mg [10], Al-Ca-Mg [11] and Ca-Mg-Si [12]. Earlier work was summarized previously [3].

3

Phase diagram of the Al-Ca-Mg system as basis for new alloys

The ternary Al-Mg-Ca phase equilibria were investigated using X-ray diffraction methods, metallography, scanning electron microscopy with energy and wave-length dispersive X-ray microanalysis and differential thermal analysis. Large ternary solubilities were found for the binary phases CaMg2, Al2Ca and Al3Ca8 [11]. The isothermal section at 300°C including the investigated sample compositions is shown in Figure 2. The calculated liquidus surface is given in Figure 3. In compositions near the ternary monovariant eutectic L → CaMg2 + Al2Ca remarkable microstructures formed of rod-like crystals were found, that are identified as intergrowth between CaMg2 and Al2Ca.

Ca

This work single phase two-phase three-phase

100

300°C 80

at. %

Ca

Al3Ca8 60

AlCa

40

CaMg2

Al2Ca Al4Ca

20

0

Mg

0

(Mg)

20

40

γ

ε

at.% Al

60

β

80

100

(Al)

Al

Figure 2: Calculated isothermal section of the Al-Ca-Mg system at 300°C including compositions of own samples. Dots represent the investigated samples compositions, arrows point to the identified phases. Threephase triangles are shaded.

15

Ca 100

50 μm

50 μm

Al35Ca50Mg15

Al30Ca50Mg20 80

YY

Ca 60

YY

at. %

YY

Y Y

YY

50 μm

50 μm

Al34Ca46Mg20

Al27Ca33Mg40 Max1

YY

40

YY

20

50 μm

50 μm

Al30Ca25Mg45

Mg

0

20

Y Y

0

Y Y

Y Y

Al20Ca20Mg60 60

40

80

Al

at.% Al

Figure 3: Calculated liquidus surface with optical microscopic images of the microstructure of the alloys near the eutectic line (solid circles). All alloys in as-cast condition, size marker 50 μm.

4

Thermodynamic calculation for AS alloys and modification with Ca

Figure 4 shows a calculated Mg-Al-Si vertical phase diagram section at constant 1 mass% Si through the alloys AS21 and AS41. The reliability of this calculation is verified with our DTA measurements of two key samples with the compositions of these two alloys. It may be tempting to try to improve the mechanical properties of AS alloys by the addition of Ca. Without knowledge of the phase equilibria one might expect Al2Ca as possible, strength improving precipitation. In Figure 5 the calculated effect of adding 0.1 mass% Ca to AS alloys is shown. The very stable ternary phase CaMgSi is precipitated in a wide range of alloys even at this low Ca-addition. At higher Ca-addition (0.5 mass% Ca) the phase CaMgSi is formed primary during solidification and the liquidus temperature would increase to around 700 °C. The envisaged phase Al2Ca is formed only at very high Ca-additions (more than 3 mass% Ca).

16 700

AS21

DTA

L

650

700

AS41

650

L + (Mg)

L + (Mg) L + (Mg) + CaMgSi

L + (Mg) + Mg2Si

550

T/°C

T/°C

L + CaMgSi 600

600

L + (Mg) + CaMgSi + Mg2Si 550

(Mg) + Mg2Si

(Mg) + CaMgSi + Mg2Si 500

500

1

2

Mg 99 Al 0 Si 1

3

Mg 98.9 2 Al 0.0 Ca 0.1 Si 1.0

4

Mg 94 Al 5 Si 1

wt.% Al

Figure 4: Calculated vertical section in the ternary Al-Mg-Si system at constant 1 mass% Si through the alloys AS21 and AS41. Thermal signals of key samples verify the reliability of the calculation.

4

6

8

Mg 88.9 Al 10.0 Ca 0.1 Si 1.0

wt.% Al

Figure 5: Calculated vertical section in the quaternary Al-Mg-Si-Ca system at constant 1 mass% Si and 0.1 mass% Ca. The very stable ternary phase CaMgSi is precipitated in a wide range of alloys even at this low Ca-addition.

5. Replacing RE with Ca in AE alloys Mg-alloys with RE additions are known for their good high temperature properties and corrosion resistance. Replacing RE with Ca in Mg-alloys could lower the cost, without impairing the mechanical properties. Pekguleryuz and Baril [1] reported the same level of creep resistance for AX alloys than for AE alloys at much lower cost. They observed Al2Ca as new strengthening phase in their samples. The calculated phase fractions during solidification of the alloy AX51 is given in Figure 6. Four different equilibria can be observed. Al2Ca is formed below 500 °C as dominant precipitate. At lower temperatures additional γ-Al12Mg17 may be formed. Figure 7 compares the liquid phase fraction of AE42 during equilibrium solidification (lever rule) with the case of blocked solid state diffusion (Scheil model). The large difference indicates the potential for segregation of this alloy, especially during secondary solidification. Cerium was used as prototype RE for this calculation. 1.0

1

(Mg)

3

4

2

1

Liquid 0.8

0.1

γ-Al12Mg17

f(Liquid)

Phase fractions, f

φ

"AX51" 1 L + (Mg) 2 L + (Mg) + Al2Ca 3 (Mg) + Al2Ca 4 (Mg) + Al2Ca + γ Al2Ca 200

300

400

0.4

1

lever rule (Mg)

0.01 100

0.6

"AE42" 1 L + (Mg) 2 L + (Mg) + Al2Ce 3 L + (Mg) + Al11Ce3 4 L + (Mg) + Al11Ce3 + γ-Al12Mg17

500

600

700

T/°C

Figure 6: Calculated molar phase fractions during solidification of the alloy MgAl5Ca1(mass%), "AX51".

0.2

Scheil model 3

4 0.0 400

450

500

550

2 600

650

700

T/°C

Figure 7: Calculated fraction liquid during solidification of Mg-Al4Ce2 (mass%), "AE42", lever rule (light line) and Scheil model (solid line).

17 An additional application of thermodynamic calculation is calculation of enthalpies of real alloys. An example is shown in Figure 8. Here enthalpy of solidification is given for the alloys “AE42” and “AX51” as specified above. The liquidus temperatures are quite similar, but the temperature interval of solidification is substantially wider for the Ca-containing alloy AX51 than for AE42. This type of calculation can also be performed for solidification under Scheil conditions using the present thermodynamic Mg-alloy database. 30

H / kJ/mol

25

620°C

AX51 20

15 514°C

10 500

625°C

AE42

580°C

550

600

650

700

T/°C Figure 8: Calculated enthalpy during equilibrium solidification for alloys AE42 and AX51. Liquidus and solidus temperatures are given for each alloy.

References [1] M. O. Pekguleryuz, E. Baril, Mater. Trans. 2001, 42, 1258-1267. [2] F. v. Buch, S. Schumann, H. Friedrich, E. Aghion, B. Bronfin, Metall 2002, 56, 40-48. [3] R. Schmid-Fetzer, J. Gröbner, Adv. Eng. Mater. 2001, 3, 947-961. [4] S.-L. Chen, S. Daniel, F. Zhang, Y. A. Chang, W. A. Oates, R. Schmid-Fetzer: On the Calculation of Multicomponent Stable Phase Diagrams. J. Phase Equilibria, 2001, 22, 373-378. [5] D. Kevorkov, R. Schmid-Fetzer, Z. Metallkd. 2001, 92, 946-952. [6] D. Kevorkov, R. Schmid-Fetzer, A. Pisch, F. Hodaj, C. Colinet, Z. Metallkd. 2001, 92, 953-958. [7] D. Kevorkov, J. Gröbner, R. Schmid-Fetzer, J. Solid State Chem. 2001, 156, 500505. [8] J. Gröbner, D. Kevorkov, R. Schmid-Fetzer, J. Solid State Chem. 2001, 156, 500505. [9] J. Gröbner, R. Schmid-Fetzer, A. Pisch, C. Colinet, V. V. Pavlyuk, G. S. Dmytriv, D. G. Kevorkov, O. I. Bodak, Thermochimica Acta 2002, 389, 85-94. [10] J. Gröbner, R. Schmid-Fetzer, Intermetallics 2002, 10, 415-422. [11] J. Gröbner, D. Kevorkov, I. Chumak, R. Schmid-Fetzer, Z. Metallkd. 2003 (in press). [12] J. Gröbner, I. Chumak, R. Schmid-Fetzer, 2003 (submitted for publication).

18

Development of the new Creep Resistant Alloy AS31 Vadim V. Agalakov, JSC “AVISMA” , Berezniki, Russia

Introduction Lately there is no lack of research activities connected with the development of new magnesium alloys with improved creep resistance. The possibility to substitute aluminum for magnesium in engine and transmission components is more than tempting. The selection of a new alloy is based on several main principles which are well-known. These are: adequate mechanical properties, good castability, satisfactory corrosion resistance, a relatively low price, ecological and economic aspects of scrap recycling, good machinability. Mg-Al-Zn-Mn and Mg-Al-Zn alloy families can meet all these requirements but their principle weakness is in inadequate creep resistance. The present approach to overcome this disadvantage is to alloy with rare-earth elements and strontium [1, 2, 3], but it is evidently that it will hardly allow to develop a reasonably priced alloy though having excellent mechanical properties and improved creep resistance. There is also an Mg-Al-Si alloys family [4, 5, 6]. It was not long ago when these alloys were put on the list of commercial alloys, though they have been used by Wolkswagen since 1970s. These alloys are more competitive in price than those alloyed with rare-earth metals. They have better elevated temperatures properties than such conventional alloys as AZ and AM. The alloy AS41 was developed in response to the requirement to improve creep resistance at temperatures up to 150 °C. The alloy was used in crankcases. High strength of this alloy is attained due to the stable phase Mg2Si which deposits on grain boundaries. The alloy AS21 has better creep resistance than that of AS41 but lower strength at the room temperature and worse castability. We decided to develop a Mg-Al-Si based alloy (table 1). In fact we had to settle three problems: to obtain good results on corrosion resistance and castability and maintain satisfactory mechanical properties.

Chemical composition of AS31 Chemical composition of AS31 is given in Table 1. Table 1. Chemical composition of AS31

Alloy

Basic components in % wt.

Mg Al Si Mn AS31 Matrix 2.5- 0.8- 0.243.4 1.1 0.34

Impurity content in % wt., max Zn 0.110.25

Be 0.00050.0015

Cu 0.008

Ni 0.0010

Fe 0.004

Other 0.01

Increased creep resistance of AS31 is explained by the formation of the high-melting intermetallide Mg2Si ( Figure 1. ) and the decreased content of the low-melting intermetallide Mg17Al12 being part of the eutecticum.

19 Mg2Si Intermetallic (Chinese script)

Figure 1. Mg2Si Intermetallic ( Chinese script )

AS31 ingots microstructure Microstructure of upper and lower parts of an 8-kg ingot is shown in Figures 2 - 3. Microstructure of both ingots are similar : Mg-base with the thermodinamical stable Mg2Siphase (‘Chinese script’) and Mg17Al12-phase in the eutecticum, consisting of lamellar white Mg, Mg17Al12 dark plates and globular Mg12Al17. All phases are visible on grain boundaries. In the gate part there is a small micro-porosity zone with little precipitated needles. The most of the ingot surface has neither micro-porosity nor needles. The ingots of both sizes have no gas pores. Microstructure satisfies customer requirements.

Figure 2. AS31 Alloy microstructure ( the top of the 8-kg ingot)

20

Figure 3. AS31 Alloy microstructure (the bottom of the 8-kg ingot)

Conclusion The ingots of both sizes are of good quality and homogeneity. The surface quality of the 8-kg ingot is better than that of the12-kg ingot. The latter also totally satisfies customer requirements.

Mechanical properties of AS31 Tensile samples were die cast and tested at JSC ‘VSMPO’ in accordance with ASTM B557 guidelines. The test results are presented in Figure 4. Similar mechanical properties data were obtained within USCAR MPCC Program.

AS31 creep properties An AS31 creep resistance investigation was also carried out at JSC ‘VSMPO’ using die cast samples. Samples were tested at 150oC during 100 hours under different loads (Figure 5). The chemical composition of samples is Al – 2.95%, Si – 0.94%, Mn – 0.26%, Zn – 0.18%.

21 Ultimate Tensile Strength

RT

250,0

125

Sress (MPa)

200,0

150 180

150,0

100,0

50,0

0,0

AS31

Stress (MPa)

Tensile Yield Strength 140,0

RT

120,0

125

100,0

150 180

80,0

60,0

40,0

20,0

0,0

AS31

% Elongation

RT

20,0

125

18,0 16,0

150

% Elongation

14,0

180

12,0 10,0 8,0 6,0 4,0 2,0 0,0

AS31

Figure 4

22 Compression Yield Strength

RT

100,0 90,0

125

80,0

150

Stress (MPa)

70,0 60,0 50,0 40,0 30,0 20,0 10,0 0,0

AS31

Figure 4. Mechanical tests results Tensile Creep, 35-70 MPa - AS31 Tensile Creep at 100 Hours - 35-65 MPa Tensile Creep at 500 Hours - 70 MPa

0,80 0,70 0,60

Creep, ε %

0,50 0,40 0,30 0,20 0,10 0,00 30

35

40

45

50

55

60

65

70

75

Stress, MPa

Figure 5. Creep test results

Casting properties of AS31 Casting properties ( castability, hot brittleness) of AS31 were investigated by a customer when producing a trial lot of transmission housings. Based on the test results casting properties of AVISMA alloy AS31 were highly evaluated by an automobile company.

Corrosion tests results Comparative corrosion tests were carried out at the room temperature and complete immersion. Test results were evaluated according to GOST 9.020. Average corrosion rates of various samples tested for 24 hours in 3% NaCl are represented in Figure 6. Corrosion test results according to ASTM B-117 are represented in Figure 7.

23 1,80

157

Aver age cor r osion r ate, g/m2*h

1,60 1,40

AS

ZAC

155

AZ

1,20

156

1,00

159

149

0,80

154

148

147

165

0,60 0,40 0,20 0,00

ingot #

Figure 6. General corrosion test results of chill cast magnesium alloys at complete immersion into 3% NaCl

Alloy Comparison 10,00 AS31

9,00

AZ91D

Corrosion Rate (mils/yr.)

8,00 7,00 6,00 5,00 4,00 3,00 2,00 1,00 0,00 0

2

4

6

Test Duration (days)

Figure 7. Corrosion test results as per ASTM B-117

8

10

12

24

Summary Avisma specialists together with VSMPO specialists managed to develop the new magnesium die casting alloy AS31 having increased creep resistance at 150 °C and designated for transmission housing applications. The developed technology allows to obtain a high quality alloy totally satisfying customer requirements. The alloy was tested successfully and taken for production by several European die casters.

References [1] John F. King, Paul Lyon, Ken Savage «Influence of Rear Earth Elements and Minor Additions on Properties and Performance of Magnesium-Yttrium Alloys in Critical Aerospace Applications», Proceedings of the 59th Annual World Magnesium Conference, Monreal, Canada, 2002, p 15. [2] F. von Buch, S. Schumann, H. Friedrich, E. Aghion, B. Bronfin and others, «New Die-Casting Alloy MRI153 for Power-Train Applications», «Magnesium industry», # 8 September 2002, p.34. [3] J. F King «Development of Practical High Temperature Magnesium Casting Alloys», «Magnesium alloys and their Applications», edited by K.U. Kainer, WILEY-VCH 2000, 14-22. [4] K. Pettersen, H. Westengen, J. Ivar Skar, M. Videm and others «Creep Resistant Mg Alloy Development», «Magnesium Alloys and their Applications», edited by K.U. Kainer, WILEY-VCH 2000, 29-34. [5] P. Zhang, R. Agamennone, W. Blum and others «Microstructure and Creep Properties of Die-cast Mg-Al-base Alloys AZ91 and AS21», «Magnesium Alloys and their Applications», edited by K.U. Kainer, WILEY-VCH 2000, 716-721. [6] William E. Mercer II, «Magnesium die cast alloys for elevated temperature applications», «Magnesium in automobiles»,Ten SAE Technical Papers on Advantages of Magnesium Use, SAE July 1992, 45-53.

25

Phase Formation, Precipitation and Strengthening Mechanisms in Mg-Zn-Sn Alloys S. Cohen1, G.R. Goren-Muginstein1, S. Avraham1, G. Dehm2 and M. Bamberger1 1. Department of Material Eng., Technion IIT, Haifa 3200 ISRAEL 2. Max-Planck-Institut für Metallforschung, Stuttgart, GERMANY

1

Introduction

In recent years there has been a major effort to produce new magnesium alloys that exhibit good mechanical properties, stable microstructure at high temperatures and good castability. The most widely used Mg alloy is AZ91, in which Al is the major alloying element. Studies on the behavior of AZ91 at elevated temperatures reveal a major creep problem above 1200C. The cause of this phenomenon is the grain boundary phase Al12Mg17 (γ phase). The γ phase softens at high temperatures [1, 2]. AS41 and AS21 have better high temperature behaviors, due to the absence of γ-phase and the grain boundary pinning by the Mg2Si precipitates [3, 4]. However both these alloys have very poor castability and thus cannot be applied in the automotive industry. Various studies have been made in order to find sensational elements in order to replace part of the Al (with Zn and Ca [5-7], Y, Gd [8], or other rare earth elements [9]) or to add additional elements to improve the creep resistance in Mg-Al alloys [3, 4, 10]. Previous research on Mg-Zn alloys show that Zn improves the castability and has several stable inter-metallics with Mg [5], and that Sn forms a stable Mg2Sn compound [11]. Therefore it was decided to focus on finding an optimal MgZn-Sn composition that will have the best combination of mechanical properties and structural stability at elevated temperatures and will allow heat treatments.

2

Experimental

Three different compositions of Mg-Zn-Sn alloys, marked I, II, II, were initially studied, (see table 1). 99.9% pure Mg (DSM) was melted in a cemented-graphite crucible, then pure Zn 99.5% (Aldrich) and Sn 99.8% (Spectra) were added. A mixture of CO2 + 0.6% SF6 was used as a protective atmosphere. The melt was held at 720 °C for 20 min, followed by casting into a steel mold, forming a disc 60-mm in diameter and 9-mm thick. Table 1: The alloys’ chemical composition

Alloy

Composition [%wt]* Zn Sn Mg

I II III

2.25 4.46 4.02

3.46 3.75 9.86

Bal. Bal. Bal.

*Analyzed by MRI –Magnesium Research Institute

26 All three alloys were solution-treated using the following procedure: 96 h at 300 °C, heating to 465 °C at a rate of 1 °C/h and keeping at this temperature for an additional 96h, followed by water quenching. Aging treatment was conducted in molten salt (sodium nitrate 50%+ potassium nitrate 50%) in the range of 150-275 °C. Vickers micro hardness measurements, using an Akashi MVK-E3 apparatus with a load of 100 grf, served to trace the precipitation hardening of the alloys. The specimens were sliced at 1/3 of their height for microstructural characterization, thereby ensuring the same solidification conditions for all samples. The samples were polished and then etched, the as-cast specimens with a solution of 4 vol.% nitric acid + ethyl alcohol, and the heat-treated specimens with a solution of 5 gr picric acid + 5 gr acetic acid in 100 ml ethyl alcohol. The specimens were prepared for SEM, EDS, Auger spectroscopy and XRD inspection. The thermo-dynamic simulations were done by using the Thermo-Calc software.

3

Results and discussion

3.1

Choosing the alloy composition

An isothermal section at 450 °C of Mg-Zn-Sn system, seen in fig. 1, was calculated using the Thermo-Calc software, in order to determine the most suitable composition for solution treatment. The solid solution α-Mg has solution boundaries at approximately 5% Sn and 5% Zn. Higher percentage of Sn will result in a precipitation of Mg2Sn, and a higher percentage of Zn will result in residual liquid in the grain boundary. Thus it was decided that the Zn content will be in the middle of the α-Mg single phase region at the first alloy and to work at the edges of the second and the third alloy at 4.5 wt%. Two compositions in the α-Mg single-phase region (Sn content below 4.5 wt%) and one with primary precipitation of Mg2Sn at 9 wt%, as presented in table 1 were chosen. The phases as a function of the temperature for the three alloys are presented in fig 2 a, b, and c. The calculated diagrams show the existence of four phases 1) liquid – that will transform into a eutectic structure, 2) α-Mg, 3) MgZn and 4) Mg2Sn.

Figure 1: Mg-Sn-Zn Thermo-Calc calculated phase diagram at 450 °C.

All three as cast alloys were examined by XRD as shown in fig 3 a, b, c, and SEM micrographs, shown in fig 4 a-c. The expected α-Mg, MgZn and Mg2Sn are clearly discernible.

27

Figure 2: Thermo-Calc calculated phase diagrams for; a) alloy I, b) alloy II, and c) alloy III

a

b

c

alpha Mg Mg2Sn MgZn

20

25

30

35

40

45

50

2_ Figure 3: XRD spectra for the as cast condition; a) alloy I, b) alloy II, and c) alloy III

Figure 4: SEM micrographs for the as cast condition; a) alloy I, b) alloy II, and c) alloy III

XRD spectra of the solution treated alloys is seen in fig 5 a, b, and c. Alloys I and II exhibit a complete dissolution of the alloying elements; whereas in alloy III, the Mg2Sn still exists. This alloy was not suitable for this study and thus alloy II that has higher content of the alloying elements was chosen for further in-depth examination.

28 a

b

c

alpha Mg Mg2Sn

20

25

30

35

40

45

50

2_ Figure 5: XRD spectra for the solution treated; a) alloy I, b) alloy II, and c) alloy III

3.2

Precipitation

The precipitation hardening tests were made in order to track the precipitation mechanisms in the material. Several solution treated specimens were exposed to 175, 200, 225, or 250 °C, for 1-96 hrs. The hardness variation, during aging, at each temperature is shown in fig 6. Aging at 175 °C shows two hardness maxima; the first after 2 hours, and the second after 16 hours. At 200 °C the first maximum occurs after 0.5 hours and the second after 4 hours. In aging at 225 and 250 °C only a single maximum curve at about 2 hours could be observed. Micro Vickers hardness (HV)

90 85

175c 200c

80

225c 250c

75 70 65 60 55 50 0

10

20

30

40

50

60

70

80

90

100

Aging Time [h]

Figure 6: Precipitation hardening graphs as a function of time at: 175 °C, 200 °C, 225 °C and 250 °C

XRD spectrums for aging at 175 °C for 4, 16, and 48 hours, are shown in Fig 7 a-c, respectively. After the first 4 hours only α-Mg was detected, after 16 hours MgZn appeared, and after 48 hours, Mg2Sn appeared. It seams that the first pick corresponds with the precipitation of MgZn and the second pick corresponds with the Mg2Sn. In order to confirm the findings, Auger spectroscopy was used, and a ratio of 2:1 Mg-Zn was found. Mg2Zn or Mg7Zn3 are a meta-stable phase that only exists in a very narrow range of temperatures in the binary Mg-Zn phase diagram, but since casting is not an equilibrium process, metastable phases can occur.

29 a

b

c

alpha Mg Mg2Sn MgZn

20

25

30

35

40

45

50

2_ Figure 7: XRD spectra for the aged alloy II at 175 °C, 4 hrs., 16 hrs., and 48 hrs. M g Z nSn A llo y II T her mal Exp o sur e 2 0 0 c 80 I I T hEx 2 0 0 c

75 70 65 60 55 50 0

5

10

15

20

25

T ime [ d ay]

Figure 8: Hardness variation of thermally exposed specimen for 33 day at 200 °C

Figure 9: Thermally exposed specimen for 33 day at 200 °C

30

35

30 3.3

The as-cast alloy at elevate temperature

Samples cut from alloy II in the as-cast state were exposed to 200 °C for very long periods (1-33 days). The hardness variation with time is shown in Fig 8. The hardness increases in the first 24 hours from 60 HV to 70 HV and then decreases very slowly with a time of up to 65 HV. No measurable changes were detected by XRD during the thermal exposure. A SEM micrograph of the specimens exposed to 200 °C for 33 days is shown in Fig. 9. Very fine precipitates are located near the grain boundary, in a ribbon like shape of about 10 μm wide. The precipitates are 200-500 nm long and have an oval shape. Additional work must be done in order to find their chemical composition.

4

Summary and conclusions

• Alloy II was found to be the most suitable for solution treatment and aging. • The alloy has two maxima of hardness during aging. The first is associated with MgZn and the other is associated with the precipitation of Mg2Sn. • Maximum hardness of 75 HV was reached by aging at 175 °C, but over-aging occurs very rapidly and thus the alloy is unsuitable for precipitation hardening. • An additional Mg-Zn phase has precipitated during the aging, the meta-stable Mg2Zn. • Exposing the as-cast alloy to elevated temperature reveals a relative stability with time. There is a rapid increase in hardness at the beginning due to precipitation near the grain boundaries that is followed by a very slow decrease is hardness with time, due to microstructure coarsening.

5

Acknowledgement

The study was partially supported by the German Israeli foundation for scientific research and development (GIF) under contract number I-704-43.10/2001. The foundation is acknowledged for its support.

6

Bibliography

[1] M. Regev, A. Rosen, and M. Bamberger, Met. Trans., 2001, 32A, 1335-1345 [2] W. Blum, B. Watzinger, and P. Weidinger, Proc. Of the Mg Alloys and their Application Conf. (B.L. Mordike and K.U. Kainer eds.) Werksoft-Informationsgesellschaft mbH, Wolfsburg, Gwemany, 1998, p. 49-60 [3] J.J. Kim, D.H. Kim, K.S. Shin and N.J. Kim, Scripta meter.,1999, 41, 333-340 [4] G.Y. Yaun, Z.L. Liu, Q.D. Wang, and W.J. Ding, Mater. Letters, 2002, 56, 53-58 [5] A.A. Lou nad B.R. Powel in Mg Technology 2001 (J. Hryn Editor) Proc. Of the TMS symposium, 2001, P. 137-144 [6] J.F. Nie and B.C. Muddle, Scripta Meter., 1997, 37, 1475-1481 [7] A. Finkel, L. Shepeleva, M. Bamberger and E. rabkin, in Mg Technology 2002, (H.I. Kaplan Editor) Proc. Of the TMS symposium, 2002, p. 81-86 [8] L.L. Rokhlin , N.I. Nikitina, J. of Alloys and Compounds, 1998, 279, 166–170 [9] Y. Lu, Q. Wang, X. Zeng, W. Ding, C. Zhai, and Y. Zhu, Mater. Sci. and Eng., 2000, A278, 66–76 [10] Y. Guangyin, S. Yangshan, and D. Wenjiang, Scripta mater., 2000, 43, 1009–1013 [11] M. Mabuci and K. Higashi, Acta Meter., 1996, 44, 4611-4618.

31

Die Casting for High Performance – Focus on Alloy Development Dr. Per Bakke, Dr. Håkon Westengen Hydro Aluminium a.s., Magnesium Competence Centre, Porsgrunn, Norway

1

Introduction

By selecting the right combination of alloy and processing method, a wide range of temperature exposed drive train parts can be made out of die cast magnesium. A schematic illustration of the creep/strength performance vs. temperature requirements is shown in Fig. 1, indicating that the engine blocks and the automatic transmissions may be the most demanding components. Hydro Magnesium has approached the elevated temperature market from a pragmatic angle, realizing that successful new alloys for this purpose must fulfil a multitude of requirements to offer a viable solution [1-5]. This is illustrated in Figure 2. Since so many factors need to be addressed, and each end user application has its own requirements, the focus is to select alloys based on the customers’ requirements and at the same time optimise other factors. The Property Profile

Creep/Strength performance

Corrosion

Manual transmission 4WD transfer case

Intake manifold Cylinder head cover

Aut. transmission Low power

Engine blocks Aut. transmission High power

Manual transmission High power Engine brackets Oil pans, non-structural

Molten Metal Handling

Bed plate Oil pans, structural

125

150

Mechanical

Die Casting Operation

Casting Defects

The Alloy Production & Recycling Profile Ingot Casting

Reinforced engine block

100

Physical

The Diecasting Performance Profile

Remelting Recycling

Alloying Processes

The Feasibility Profile

175

Temperature (oC)

Figure 1. Creep strength/performace vs. temperature requirements for some automotive drivetrain components.

HES Sustainability

Ingot Cost

Raw Material Cost

Figure 2. Elements in the alloy development process.

Based on extensive R&D efforts, Hydro Magnesium is currently focusing on alloys of the AE family, which will be highlighted in this paper. The letter E denotes rare earth elements or misch metal, in this case a standard grade Ce based alloy (>50% Ce) with La (20-35%), Nd (10-20%) and Pr (4-10%). The AE alloys show a good overall performance related to the factors described in Figure 2. This includes alloy ingot production and recycling characteristics, as well as raw material and ingot cost. Die casting performance depends upon the alloy composition necessary to fulfil the high temperature and creep requirements. The key challenge is to identify alloys within this family, which fulfil the property requirements, and at the same time demonstrate a die castability, which is compatible with the geometry of the targeted components [4-6].

32

2

Experimental

Test bars and plates of different experimental AE alloys were die cast on a 420 ton Bühler SC42D Evolution machine, with a five cavity plate/bar die (version 0), Fig. 3. Specimens of reference die casting alloys AZ91D, ASE210 (AS21X), AE42, AJ52 and AXJ530 were prepared and cast in the same die. Permanent mould cast medallions were withdrawn from each alloy and analysed by ICP-AES. Standard 6 mm test bars from each composition were tensile tested in a 10 kN Instron machine at temperatures ranging from room temperature and up to 210 °C in accordance with ISO 6892. In all tests a strain rate of 1.5 mm/min was used until 0.5% strain, whereafter the strain rate was increased to 10 mm/min. Two specimens from a series of alloys were tensile creep tested for minimum 150 hrs in accordance with ASTM E139, at Figure 3. The five cavity plate/bar die for the Bühler different realistic loads and temperatures. SC42D Evolution die casting machine. Tensile creep testing at 175 °C/40 MPa and 175 °C/75 MPa were also performed on samples that were aged at 175 °C for 500 hrs. This because phase transformations possibly taking place during prolonged exposure at elevated temperatures could negatively affect the creep properties. This information would be highly relevant for e.g. engine blocks, transmission housings or bed plates in service. Stress relaxation testing was done at 175 °C in accordance with ASTM E328-86. Initial loads were 12 kN and 6 kN corresponding to 106 and 53 MPa, respectively. The specimens (12 mm dia., 5 mm length), were cut from one arbitrary end of one arbitrary tensile creep bar. Die cast 3 mm test plates were subjected to corrosion testing in accordance with ASTM B117 (salt spray test). Sample preparation included vibratory finishing (90 min) and immersion in cold chromic acid, 180 g/l. Reference plates of AZ91D, AM60B and AE42 (5 plates each) were tested in the same chamber.

3

Results and discussions

3.1

Tensile testing

Tensile yield strength (TYS) as a function of temperature for some experimental AE alloys is shown in Fig. 4. AE42 and aluminium alloy A380 are shown for comparison. It is seen that TYS for the experimental AE alloys is well above that of AE42. Specifically, at 200 °C, TYS for four of the AE alloys is in the range of 105-115 MPa. For AE42, TYS at 200 °C is measured to 90 MPa. On the other hand, TYS for A380 at 200oC is as high as 175 MPa. Ultimate tensile strength (UTS) for the same alloys is shown in Fig. 5. It is seen that at room temperature all AE alloys tested exhibit UTS in the 225-250 MPa range. The experimental AE alloys shown perform better than AE42 at all temperatures. At 200 °C, UTS of the AE alloys is in the range of 120-135 MPa, while for A380 it is 240 MPa.

33 325

195

300

180

150 135

250 UTS (MPa)

165 TYS (MPa)

275

A380 AE v7 AE v8 AE v9 AE v10 AE v2 AE42 .

200

A380 AE v7 AE v8 AE v9 AE v10 AE v2 AE42

175

120

150

105

125

90

100

0

50

100

150

200

250

o

Temperature ( C)

Figure 4. Tensile yield strength as a function of temperature for some AE alloys, AE42 and A380. 35

225

0

50

.

100

150

200

250

o

Temperature ( C)

Figure 5. Ultimate tensile strength as a function of temperature for some AE alloys, AE42 and A380.

Elongation at fracture (%)

The range of alloying elements in AE type alloys is quite wide due to the extensive liquid solubility [7]. Compared to other disper25 soid forming alloying elements like Sr and 20 Ca, RE elements have a less harmful effect 15 on ductility/fracture toughness. Elongation at . 10 fracture as function of temperature is shown in Fig. 6. It is seen that the ductility of the AE 5 alloys is better than the reference alloys at 0 room temperature as well as at the higher tem0 50 100 150 200 250 peratures. Temperature ( C) It should also be noted that the value of the Figure 6. Elongation at fracture as a function of yield stress is very dependent on the strain temperature for some AE alloys, AS21X, AE42 rate, especially at elevated temperatures, and and A380. direct comparison of values from different sources may be misleading. Therefore we have chosen to use only our own data in Figs. 4-6. 30

AE v7 AE v8 AE v9 AE v10 AE42 AS21X A380

o

3.2

Tensile creep

In Figs. 7-8 creep strain as a function of time are shown for selected alloys at 150 °C/90 MPa and 175 °C/75 MPa, respectively. It is seen that at both these relatively demanding conditions the creep curves of some of the new AE alloys are significantly better than for example AE42. The new AE alloys are matching AJ52, which is known for outstanding creep performance up to 175 °C (but poor die castability) [8]. For applications exposed to such conditions AS21X is no longer an option, although it may be an exceptionally good solution for modest temperature/load range applications. Total creep strain at 100 hrs seems to be better for the new AE alloys than for A380. However, the creep rate of A380 beyond 100 hrs is likely to be as low as the best Mg alloys shown in Figs. 7 and 8. Creep rates at 100 hrs are shown in Figs. 9-10. The results are based on the average of two parallels between 80-120 hrs. The resolution of the tensile creep measurements is at -7 -1 the level of 5 ⋅ 10 hr . From Fig. 9 it is seen that new AE alloys at 175 °C perform similar to AE42 up to 50 MPa. At higher loads the new AE alloys show substantially lower creep rates. AS21X may not be considered for applications at 175 °C.

34 10

AZ91 AS21X AE42 A380 AE v8 AE v9 AJ52

1

Creep strain (%)

Creeep strain (%)

10

0.1

0.01 1

10

100

AS21X AE42 A380 AE v8 AE v9 AE v7 AJ52

1

0.1

0.01

1000

1

10

Time (hrs)

100

1.00E-03

AS21X AE42 AJ52 AE v4 AE v7 AE v8

1.00E-04 1.00E-05 1.00E-06

Figure 8. Creep strain as a function of time at o 170 C/75 MPa for selected alloys. Notice that both axes are exponential. 1.00E-04 Creep rate (hrs-1)

Figure 7. Creep strain as a function of time at 150 oC/90 MPa for selected alloys. Notice that both axes are exponential.

Creep rate (hrs-1)

100 Time (hrs)

As cast 1.00E-05

Agedtreated Heat

1.00E-06

1.00E-07

AE-v2 AE-v6 AE-v4 AE-v1 AE42 AS21X AE-v7 AE-v8

1.00E-07 30

40

50

60

70

80

90

Load (MPa)

Figure 9. Creep rates at 100 hrs as a function of load at 175 oC for AS21X, AJ52 and some AE alloys.

Figure 10. Creep rates at 100 hrs at 175 C/40 MPa for some as cast and aged AE alloys, with reference alloys AE42 and AS21X. o

Specimens were aged for 500 hrs at 175 °C. From Fig.10 it is seen that the creep rates at 175 °C/40 MPa of the as cast and aged samples are almost indistinguishable. An exception is AE42 where the creep rates of the aged specimens are seven times higher. The reason is that during solidification, the AE alloys with a RE/Al ratio above a certain critical level form Al2RE particles rather than the less stable Al11RE3. During aging the Al11RE3 particles (as in AE42) transform into Al2RE and Al, which subsequently form the brittle Mg17Al12 phase, leading to a lower creep resistance [9]. This indicates that the improved AE alloys in service are expected to exhibit the same creep properties regardless of operating time. 3.3

Stress relaxation

In the stress relaxation test a compressive initial load is applied, and as the alloy deforms due to creep, the remaining load is measured. Bolted joints experience the same type of behaviour in components [10]. Results of the stress relaxation tests are shown in Fig.11. At 12 kN initial load there are AE alloys outperforming the reference alloys AJ52 and AE42. Specifically, at 100 hrs, the alloy AE v7 has >70% initial load remaining. The internal ranking of the alloys tested is the same at 6 and 12 kN, as well as at 150 °C, 12 kN. Notice that the tensile yield strength is an absolute upper limit for the load that can be applied in a creep test. The initial load of the stress relaxation test (106 MPa at 175 °C) is well above the tensile yield strength of for example AS21X. This means that the results of AS21X at 106 MPa/175 °C are misleading as remaining load quickly drops to below TYS.

35 100 90

Remaining load (%)

80 70 60 50 40

AE v7 AE v1 AE v2 AE v8 AJ52 AE42 AS21X

30 20 10 0 0

25

50

75

100

125

150

175

Time (hrs)

Figure 12. Corrosion rate (MCD = mg/cm2/day) of AE alloys as a function of Al and RE contents.

Figure 11. Remaining load as a function of time for some AE alloys at 175 °C with initial load 12 kN (106 MPa). The alloys A380, AS21X, AE42 and AJ52 are included for reference.

3.4

Corrosion and other requirements

Generally, corrosion does not represent any problems with AE alloys [11]. Fig. 12 shows results of the corrosion testing. MCD values (mg/cm2/day) are plotted as function of Al and RE contents. The corrosion rates are low and consistent, in the range of 0.0240.069 mg/cm2/day within the composition range shown. It is seen that the corrosion rates decrease slightly with increasing Al content. Results of reference samples AZ91D, AM60B and AE42 tested alongside the AE alloys confirm the low corrosion rates of the AE alloys. Corrosion values at this level can of course only be obtained as long as the contents of the heavy metal impurities Fe, Ni and Cu are kept below the maximum limits specified in the High Purity standards; 44, 14 and 84 ppm, respectively. The proposed AE alloys also fulfil the requirements with respect to recyclability, compositional stability and traditional melt handling. The AE alloys cannot only be recycled in flux-free processes, but also with traditional flux. This is illustrated in Fig. 13 where ΔG of formation of some chlorides is shown [12]. It is seen that CaCl2 and SrCl2 both are significantly more stable than MgCl2, hence Ca and Sr in molten alloys (AX and AJ al-300 loys) in contact with MgCl2 conMnCl taining salts will react with the salt -400 forming CaCl2 and SrCl2 on the MgCl expense of MgCl2. The chlorides of -500 / NdCl / CeCl the RE elements are slightly more / LaCl -600 stable than MgCl2, but RE in alloys 2 NaCl CaCl SrCl in contact with MgCl2 containing -700 BaCl 2 KCl salts react very slowly. Thus pro-800 duction and recycling processes for 500 600 700 800 900 AE alloys can be run with tradiTe mpe rature (C) tional salts of the MgCl2-CaCl2Figure 13. ΔG of formation as a function of temperature for NaCl-KCl-CaF system. 2 ∠ G (kJ/mole)

2

2

2

3

3

2

3

3

2

2

3

3

2

2

some common chlorides. Thermodynamic data from [12].

36 Normal gas protection systems, either by diluted SO2 or SF6, works exactly as with the traditional alloys. Also, there is no reason to expect larger dross or sludge accumulation than for instance with AM50. Freezing during prolonged stoppages and subsequent remelting is straightforward and yields no negative surprises. Physical properties, thermal data, and die castability will be reported in upcoming papers.

4

Conclusions

By combining well-known principles of alloy design, with the characteristics of the die casting process, magnesium alloys can be tailored to specific property requirements. The tensile-, creep- and corrosion results demonstrate that AE alloys show a remarkable potential among the Mg-based alloys. If we include recyclability, the AE alloys should become a very attractive light weight solution for the automotive industry.

References [1] H. Westengen, P. Bakke, SINOMAG Conf., Beijing, China, Nov. 2002 [2] H. Westengen, P. Bakke, PSTAM, Osaka, Japan, Jan. 2003 [3] P. Bakke, H. Westengen, K. Pettersen, D. Albright, SAE Technical Paper 2003-010189, 2003 [4] P. Bakke, H. Westengen, K. Pettersen, Magnesium Technology 2003, (Ed.: H.I. Kaplan), TMS-AIME, 2003. 171-176. [5] P. Bakke, K. Pettersen, H. Westengen, 10th Mg Automotive and End User Seminar Aalen, Germany, Sept. 2002. [6] H. Gjestland, S. Sannes, H. Westengen, 10th Mg Automotive and End User Seminar Aalen, Germany, Sept. 2002. [7] K. Pettersen, P. Bakke, D. Albright, Magnesium Technology 2002, (Ed.: H.I. Kaplan), TMS-AIME, 2002. 241-246. [8] M. Pekguleryuz, E. Baril, Magnesium Technology 2001, (Ed.: J. Hryn), TMS-AIME, 2001. 119-125. [9] B. Powell, V Rezhets, M.P. Balogh, R.A. Waldo, Magnesium Technology 2001, (Ed.: J. Hryn), TMS-AIME, 2001. 175-181. [10] K. Pettersen, S. Fairchild, SAE Technical Paper 970326, 1997 [11] W.E. Mercer II, J.E. Hillos, SAE Technical Paper 920073, 1992 [12] Barin, O. Knacke, Thermodynamic properties of inorganic substances, Springer Verlag, Berlin, 1973. With Supplement, 1977

37

New Mg-Alloys for the Thixomolding™ Process Julio Aguilar, Martin Fehlbier, Andreas Bühring-Polaczek Foundry Institute, Aachen University (RWTH)

Introduction As a result of the instability of the oil prices and the improving of the comfort in the automobile construction, which result in heavier cars, it is expected that magnesium with its good mechanical properties and low density will grow in importance as a construction material. However because of the high reactivity of magnesium is the melting and casting of its alloys using the conventional casting processes just under very tight safety and protective measures possible. The ThixomoldingTM process on the contrary offers safer working conditions. The metal is heated just into the semisolid state, where it is not very reactive, away from the atmosphere gases in an extruder that is directly connected to the mold [1]. Figure 1 shows an schematic view of a Thixomolding™ machine. Thixomolding™ is based on the thixotropic characteristics of semisolid metals found by Spencer and Flemimgs in 1972 [2]. The main characteristic during the processing of thixotropic metals is the presence of a solid fraction formed of globular particles dispersed in a liquid phase. There are basically three different ways to achieve this globular microstructure. The first one is the so called thixoforming in which a specially produced prematerial with globular microstructure is heated into the semisolid state and by a determined solid fraction pressed into a mold [3]. In the rheoforming processes a melt is cooled into the semisolid state under special conditions to obtain globular solid particles within the melt and then is directly pressed into a die. The rhoforming process should have the advantage of reducing the first two steps of the thixoforming processes two just one. The best known rheoforming process is the “New Rheocasting” from UBE, Japan [4]. More complex conditions exists during the one-step process of semisolid injection molding or Thixomolding™ where feedstock particulates of a specific nature are subjected to the simultaneous influence of the heat supplied from an external source and shear force imposed by the injection screw [5, 6].

Figure 1: Schematic view of a Thixomolding™ machine.

38 Until now most of the research on the Thixomolding™ process has been done using common magnesium alloys like the AZ91 or AM60. For this reason is the development and testing of new alloys with better mechanical and casting properties interesting. The objective of this work is to analyze a series of tests that were conducted together with Husky Injection Molding Systems in Bolton, Canada using mixtures of the alloys AZ91 and MEZ under different working parameters.

Experimental The experiments were conducted on a TXM 500-M70 prototype Thixomolding™ machine from Husky with a clamping force of 500 tons. Chips of the magnesium alloys AZ91 (MgA19Zn), MEZ (MgRE3Zn) and mixtures of the two of them were used as feedstock material. Because of the very low solidification range of the alloy MEZ of about 5K it was decided not to process it alone. A mixture with AZ91 was used in order two increase the solidification range of the resulting material through the addition of Aluminum. Different series of experiments were conducted using specially designed test dies from Husky two evaluate the mechanical and flow properties of the materials. To evaluate the mechanical properties three different series of experiments were conducted using a tensile bar mold with four cavities [7]. At the beginning of each series the alloy AZ91 was used as a base of comparison for the rest of the experiments. During the first set of experiments a mixture from the chips of the alloys AZ91 and MEZ in a 2 :1 relation (two parts of AZ91 and one part of MEZ) was processed under variation of the semisolid slurry temperature. In the second set of experiments the relation from AZ91 to MEZ was changed to 1 :1 and the slurry temperature varied. The third row of experiments was conducted at a constant temperature under variation of the mixture starting with pure AZ91 and then adding MEZ to obtain 9 :1, 4 :1 and 2 :1 relations. The tensile bars were then tested on a universal tension testing machine Instron 4467, Instron Co., Canton (MA), USA at a test velocity of 20 mm/s. Ten bars from every parameter were tested in order to obtain a median. Furthermore the samples were metallographically analyzed under optical and scanning electronic microscopes (SEM) and the single phases chemically evaluated with Energy-Dispersive X-Ray analysis (EDX). Table 1 shows the nominal composition of the different chip mixtures. For the evaluation of the flow properties a spiral mold with a total length of two meters was used. In a single set of experiments, the mixture relation was changed from pure AZ91 to AZ91 :MEZ mixtures with the following compositions: 9 :1, 4 :1 and 2 :1 at a constant temperature. The relation 1 :2 was tested then at different semisolid slurry temperatures. Table 1: Nominal chemical composition of the different alloy mixtures used during this work.

Chip mixture MEZ MEZ: AZ91 (1 :1) MEZ: AZ91 (1 :2) MEZ: AZ91 (1 :4) MEZ: AZ91 (1 :9) AZ91 (1)

Experimentally evaluated

Al-Gehalt [%] 0,0 4,5 6,0 7,2 8,1 9,0 (2)

Norsk-Hydro

Zn-Gehalt [%] 0,40 0,70 0,80 0,88 0,94 1,00

Ce-Gehalt [%] 2,60 1,30 0,87 0,52 0,26 0,00

T liquidus 630-640 °C(1)

598 °C(2)

39

Results Figure 2 shows the average value of the mechanical properties of the samples produced during the first set of experiments. When compared it is possible to see that the tensile and yield strength values are higher for AZ91 as those of the AZ91-MEZ 2:1 mixture relation. However the elongation of the mixture is at a slurry temperature of 610 °C of more than 5% while for AZ91 at 593 °C just of 3,88%. It is interesting to see that the yield strength is at a value of about 138 MPa almost constant for the mixture at all tested processing temperatures. In Figure 3 it is also possible to see that the curves for the AZ91-MEZ mixture run almost together and just the strain at the breakpoint changes.

Figure 2: Comparison of the mechanical properties from tensile bars produced during the first set of experiments.

Figure 3: Stress-Strain curves of tensile bars from AZ91 and AZ91-MEZ (2:1) at different temperatures.

As expected the elongation of the material increases by lower aluminum contents. Nevertheless it was found, that by the same alloy mixture the elongation varies with varying temperature. For the alloy AZ91-MEZ (1 :2) by 593 °C the elongation reaches about 3% while at 610 °C it reaches 5.1% It can be probably explained from the metallographical analysis of the samples. At a temperature of 593 °C it can be seen that some of the chips were not completely melted during the heating of the material into the semis solid state and form clusters within the sample (Figure 4b). At a higher temperature of 610 °C the number and size of this clusters is smaller Figure 4c. Figure 4a shows as a comparison the almost perfect microstructure for the Thixomolding™ process of the AZ91 alloy at a processing temperature of 593 °C and about 30% solids. This observations are confirmed by the chemical analysis of the different phases present on a sample of the alloy mixture AZ91-MEZ (1 :2) processed at 593 °C and analyzed using EDX, Figure 5. The cluster where spectrum 1 was taken has a composition similar to that of MEZ, while the rest of the analyzed regions have higher aluminum contents with no or very small amounts of rare earths. The clusters have a thickness of more than one millimeter in some sections and interrupt the continuity of the matrix.

Figure 4: Microstructure of tensile bars thixomolded at the given temperatures

40 In Figure 6 are the mechanical properties of bars produced during the second set of experiments. The alloy AZ91 presents a tensile strength of 230 MPa and an elongation of 2.6% which is lower than that obtained during the first set of experiments. A possible reason for this behavior can be that the temperature profile in the extruder was different and as a consequence the microstructure of the material is different, changing the properties of the parts. In the case of the mixture AZ91-MEZ (1:1) at a process temperature of 610 °C is the tensile strength somewhat lower but the elongation reaches 4.7%. Interesting are the values obtained with this mixtures at a process temperature of 650 °C. At this temperature the two kinds of chips are completely melted and there is no indication of solid particles present. However the tensile strength of 247 MPa and the elongation of 8.7% are the highest of all the tests made. By some samples an elongation of up to 12% was reached. Possible reasons for this behavior is the reduced aluminum content and the continuous matrix.

Figure 5: SEM picture and EDX analysis of the chemical composition of different phases and regions observed on the alloy mixture AZ91-MEZ (1:2).

An overview of the tensile strength, the yield strength and elongation of the samples produced on the third series of tests can be seen on Figure 7. The tensile strength and the yield strength drop with increasing MEZ content in the mixture while the elongation rises. At the processing temperature of 610 °C is the alloy AZ91 completely melted and do not present any solid α-phase particles. By adding MEZ to the mixture declines the aluminum content and by a mixture relation of AZ91-MEZ of 9:1 it is possible to see some solid αphase particles. As the relation increases to 4:1 and then to 2 :1 a number of not melted MEZ chips can be observed within the microstructure of the samples, Figure 8. It is again thought that this not melted particles are responsible for the decline of the elongation by the relation 2:1 although the aluminum content is lower.

Figure 6: Comparison of the mechanical properties from tensile bars produced during the second set of experiments.

Figure 7: Comparison of the mechanical properties from tensile bars produced during the third set of experiments.

41

Figure 8: Microstructure of the tensile bars of the different alloy mixtures processed during the third series of experiments at as process temperature of 610 °C.

In order to obtain information about the flow properties of the alloy mixtures examined, a series of Thixomolding™ experiments were done on a spiral mold. The mold has two spiral cavities with an overall length of two meters and the specimens are 12 mm wide and 1,5 mm thick. The experiments were done using the same process parameters used during the first and third series of tests with the tensile bars. This should enable a coupling of the two experiments to correlate the mechanical and flow properties. The flow length of the different chip mixtures were analyzed at the same metal process temperature of 610 °C, that is for AZ91 and AZ91-MEZ in relations 9:1, 4 :1 and 2 :1. Then the mixture AZ91MEZ (2 :1) was tested at 610 °C, 604 °C, 599 °C and 593 °C. For each parameter, a series of 10 samples were produced after the machine had reached stability. In Figure 9 is a comparison of the flow length from spirals produced using different process parameters. The flow length declines by lower temperatures and higher MEZ fractions in the mixture. By lower temperatures is the solid fraction higher and the heat content of the slurry lower, so that the material should tend to solidify faster. On the other hand the increasing solid fraction could have an influence on the viscosity of the material that would affect the flow properties. By higher MEZ contents is the aluminum content in the resulting alloy lower and so the flow properties of the alloy are reduced. This phenomenon has been already observed in the thixocasting process of magnesium alloys, were the flow length increases with increasing aluminum content [8]. A very interesting aspect of these evaluations is that the flow length in all cases is quite high. By a process temperature of 610 °C is for all the mixture relations higher than one meter even though the spiral have a wall thickness of just 1,5 mm. Specially for the production of parts with thin sections and long flow distances or big areas could be this characteristic determining.

Figure 9: Flow length of spiral specimens thixomolded™ at different process parameters.

42

Conclusions The results of this work show that Thixomolding™ is suitable for the processing of chip mixtures of different alloys with the objective of getting better properties and as a result increasing the quality of the produced parts. It was possible to obtain a higher elongation while maintaining a good tensile and yield strengths. However the mixing of magnesium alloy chips with different melting points present some challenges that have still to be overcome. The presence of not melted particles within the microstructure tend to reduce the ductility and had possibly also a negative influence on the flow properties of the material. An improved control of the working parameters should help to overcome this problematic. It is very important to indicate that although the results are very promising these were just the first experiments with this material. Further experiments and trials must follow to be able to improve the properties of the parts and to establish the best process window for processing the material. As a continuation of this work, the creep and heat resistance properties of the parts produced are still to be tested. The MEZ alloy was specially designed from Magnesium Elektron (MEL), England as a die casting alloy with good creep resistance. Although during this trials it was processes allways mixed witz AZ91, the creep resistance is still of interest.

Acknowledgments The investigations set out for this report received financial support from the Ministry of Economics and Technology (BMWi) by the AiF e. V. (research project, file number 42 ZN), to whom we extend our thanks. We would also like to thank Husky Injection Molding Systems for all the time, material and technical support they provided for the execution of this work.

References [1] A. Dworog: Grundlagen des Magnesiumspritzgießens (Thixomolding™), Dissertation RWTH Aachen, Gießerei-Institut: Forschung, Entwicklung, Ergebnisse, Band 35, Shaker Verlag Aachen, Germany, 2002. [2] D. B. Spencer, R. Mehrabian, M. C. Flemings: Metall Trans. 1972; 3A: 1975. [3] M. Fehlbier, J. Aguilar, P. R. Sahm: Int. J. Cast Metals Res., 2001, 14, 71–78. [4] H. Kaufmann, D. Limburg, H. Wabusseg, P. J. Uggowitzer: Das New Rheocasting – Verfahrensprinzip und Stand der Entwicklungsarbeiten, Buchauflage: Vom Werkstoff zum Bauteilsystem, LKR-Verlag, Österreich, 2000, S. 109–121. [5] R. Carnahan, R. Decker: Thixomolding™: A Revolutionary Net Scape Forming Process, Conference: Advanced Aluminum and Magnesium alloys, Amsterdam, Netherlands 20.–22. 06. 1990, p. 531–538. [6] F. Czerwinski: Scripta Materialia 48 (2003) 327–331. [7] F. Czerwinski, A. Zielinska-Lipiec, P. J. Pinet, J. Overbeeke: Acta Mater. 49 (2001) 1225–1235. [8] M. Fehlbier: Herstellung, Charakterisierung, und Verarbeitung teilflüssiger metallischer Werkstoffe am Beispiel verschiedener Aluminium- und Magnesiumlegierungen, Dissertation RWTH Aachen, Gießerei-Institut: Forschung, Entwicklung, Ergebnisse, Band 37, Shaker Verlag Aachen, Germany, 2002.

43

Structure and Morphology of Effective Obstacles in High Performance Mg – Rare Earth Base Alloys Bohumil Smola,1), 2) Ivana Stulíková,1), 2) Jitka Pelcová,1) Barry L. Mordike2), 3) 1)

Faculty of Mathematics and Physics, Charles University, Prague 2, Zentrum für Funktionswerkstoffe gGmbH, Clausthal-Zellerfeld, 3) Institute of Materials Engineering and Technology, TU Clausthal, Clausthal-Zellrfeld 2)

1

Introduction

Excellent high performance magnesium alloys based on alloying with thorium (e.g. HK31 and HZ 32), applicable up to 350 °C, are being phased out due to their weak radioactivity. The best substitution, namely commercial alloys WE54 and WE43, owe their good mechanical properties to the morphology and thermal stability of precipitates formed by T6 treatment [1, 2]. WE type alloys can be applied in the temperature range up to about 250 °C so there is still the gap of 100 °C to be covered by new high performance magnesium alloys. Complex alloys based on alloying Mg with scandium, rare earths and manganese were developed to meet the requirements of high temperature applications [3, 4, 5, 6, 7]. Although the price of Sc rather restricts the application of these alloys, the structure and morphology of the hardening precipitates and their influence on mechanical properties show the way that should be followed in design of high performance magnesium alloys.

2

Magnesium – rare earths compounds

Magnesium – heavy rare earth (RE) alloys (e.g. WE alloys Mg-Y-Nd-Zr, Mg-Gd-Nd-Zr, Mg-Gd-Y, Mg-Gd, etc.) in peak hardening state exhibit precipitate morphology similar to that observed in Mg-Th base alloys, namely dense interlocking triangular arrangement of plate-shaped precipitates parallel to all equivalent prismatic planes of α-Mg matrix [1, 8, 9, 10]. This array not only hinders the basal slip most effectively of all possible ones [11] but also is very resistant to particle coarsening [1]. Over-saturated α-Mg matrix decomposes according to the following sequence in all of those alloys [12]: α-Mg → β’’ (hex. D019) → β’ (c-base centered orthorhombic) → β1 (fcc) → β (fcc). β’’ is metastable phase, coherent to α-Mg (a ≈ 2 * aMg, c ≈ cMg); β’ (cbco) is metastable phase, semicoherent to α-Mg (a ≈ 2 * aMg, b ≈ 8 * d( 10 10 )Mg, c ≈ cMg); β1 is metastable phase (a ≈ 0.74 nm); β is stable phase isomorphous with Mg5Gd compound (a ≈ 8 * d( 10 10 )Mg = 2.23 nm). All metastable phases were observed to precipitate in the morphology described above to reach a peak hardening. Plates were parallel to the both orders of prismatic planes dependent on the alloy and/or heat treatment [1, 2, 9, 10, 13, 12]. β’ phase precipitated only as globules at β’’ plates in alloys, where β1 phase was

44 identified in the decomposition sequence, e.g. WE54, Mg-4Y-2.25Nd-0.6Zr, Mg-7Dy2.25Nd-0.6Zr, Mg-7Gd-2.25Nd-0.6Zr annealed at 250 °C [13, 14, 12]. Stable β phase precipitates in all alloys mentioned and also in complex Mg-Sc-RE-Mn alloys as coarser plates (~ 1-2 μm dia.) parallel to { 10 10 }Mg planes. The dense triangular arrangement of β’ phase fine plates was found also in complex Mg-Sc-Gd(Y)-Mn alloys after T5 treatment and after isochronal annealing up to temperature corresponding to the minimum of electrical resistivity [4, 6, 7]. This is illustrated on Fig. 1 for MgGd10Sc0.8Mn1 alloy in T5 (200°C/72h.) condition, corresponding diffraction pattern [0001]Mg is inset. Mn2Sc particles are also imaged, although their diffraction response is not recognizable.

Figure 1. MgGd10Sc0.8Mn1 in T5, dark field image of cbco phase, [0001] of α-Mg – see inset.

Figure 2. Triangular array of β phase plates parallel to { 1010 }Mg planes in WE64 deformed at 350 °C.

The triangular arrangement of coarse β phase prismatic plates contributes also to the strength of alloys – see basal dislocations in Fig. 2 passing through this arrangement in WE64 deformed in tension at 350 °C. Different compounds were found to strengthen magnesium alloys containing RE of Cerium type (Ce, Nd), e.g. QE22 and complex Mg-Sc-Ce-Mn alloys. Mg-Ce and Mg-Nd binary alloys decompose according to the following sequence [1, 8]: α-Mg → GP zones → β’’ (hex. D019) → β’ (fcc) → β (bct). GP zones are plates parallel to ( 10 10 )Mg, β’’ is metastable phase coherent to α-Mg, β’ metastable phase is isomorphous with the β1 phase mentioned above (the same lattice parameter and morphology); β phase is incoherent to α-Mg (Mg12Ce, a = 1.033 nm, c = 0.59 nm). D019 phase was found to cause the peak hardening by annealing at 150 °C, at 200 °C over-aging occurred at concentration of 1.3wt.% RE [15]. No β’ (fcc) phase was identified there, contrary to the results on the alloys with the same RE content but in the form of misch metal [16]. Rather complicated decomposition sequence was observed in QE22 alloy [17, 1,8], but different decomposition has been found in the squeeze cast QE22 [18]. Thread-like precipitates of the stable β (bct) phase parallel to (0001)Mg plane were observed there (Fig. 3), changing their morphology to the “necklace–like” one.

45

Figure 3. Thread-like β phase in QE22 isochronally annealed up to 200 °C after T4 treatment. Left grain in [ 1010 ]Mg – see inset. [18]

Figure 4. Mn2Sc phase discs parallel to α-Mg basal plane in MnY4Sc1Mn1 crept at 350 °C and 40 MPa. Dark field weak beam (0002)Mg.

“Ribbon like” morphology of β phase, consisting of plates parallel to (0001)Mg plane stacked in [0001]Mg direction, was found in squeeze cast MgCe3Sc1Mn1 alloy [7]. It changed to the rods parallel to [0001]Mg on T5 annealing. This form was also observed in MgSc6Ce4Mn1 alloy after T5 treatment [3]. All morphological modifications exhibited the same orientation relationship to the α-Mg characteristic for this stable phase [18, 16].

3

Magnesium – scandium and manganese – rare earth compounds

Magnesium – scandium binary alloys are inferior in a creep behavior to WE43 [19], but the addition of manganese and RE improves the creep resistance considerably, especially at temperatures over 300 °C [3, 6] – see Table 1. This high creep resistance of the Sc and Mn containing magnesium alloys at high temperatures is due to the precipitation of dense Table 1. The minimum creep rates at 350°C and 30 MPa.

Alloy

Minimum creep rate [s-1]

Alloy

Minimum creep rate [s-1]

WE43 T6 MgGd15 T6 MgGd10Sc1Mn1 T5 MgGd5Sc0.3Mn1 T5 MgSc15Mn1 T5

1.10 x 10-6 4.39 x 10-7 2.70 x 10-8 2.10 x 10-8 1.84 x 10-8

MgSc6Ce4Mn1 T5 MgCe3Sc1Mn1 T5. MgY4Sc1Mn1 T5 MgY4Nd2Sc1Mn1 T5 MgY4Zn1Mn1 T5

1.33 x 10-8 1.05 x 10-8 5.42 x 10-8 7.00 x 10-9 1.62 x 10-8

46 distribution of fine Mn2Sc phase discs parallel to the basal plane of the α-Mg matrix. In Fig. 4 Mn2Sc discs are imaged in weak beam (0002)Mg reflection in MgY4Sc1Mn1 alloy crept at 350 °C and 40 MPa. Mn2Sc is the equilibrium phase of Mn-Sc system with hexagonal P63/mmc structure (a = 0.5033 nm, c = 0.8278 nm). Diameter of discs is about 10 20 nm and thickness in the middle ~5 nm after T5 treatment in all alloys. The size of discs does not change substantially on the creep exposure at 350°C. Coherent strain contrast is observed only in (0002)Mg reflection (see Fig. 4) indicating that discs strain the matrix mainly in [0001] direction. Mn2Sc discs are not very strong obstacles to the basal slip. That was proved by only moderate hardness increase as measured at room temperature in alloys, where they were the only hardening precipitates [20, 3]. On the other hand, they are very effective obstacles for cross slip of basal dislocations and for non-basal slip at higher temperatures. It follows from the effect of obstacles shape and arrangement on the flow stress needed for their overcoming [11] and manifests in very low creep minimum rates observed – Table 1. Manganese can form a variety of binary compounds with REs (inclusive Y, except La, Ce, Eu and Yb) [21]. Transition phase containing Mn and corresponding RE was observed in all Mg-RE-Sc-Mn alloys with low Sc content after T5 treatment and after isochronal annealing to the minimum of electrical resistivity [22, 4, 6, 7, 23]. It precipitated in the form of very thin plates parallel to the basal plane of α-Mg matrix – see Fig. 5. Hexagonal structure with a ≈ 2*d( 10 10 )Mg = 0.556 nm, c ≈ cMg = 0.52 nm or c ≈ 3 * cMg/2 = 0.78 nm was determined from electron diffraction patterns [7]. Recently [24] the same structure and form has been found for GP zones in Mg-2.4RE-0.4Zn-0.6Zr alloy (RE in form of Ce misch metal) after homogenization and annealing at 200°C/48h. Very thin basal plates have high aspect ratio and can be very effective obstacles to cross-slipping basal dislocations at elevated temperatures (up to ~275 °C).

Figure 5. Very thin hexagonal plates parallel to basal plane of α- Mg matrix in MgY4Sc1Mn1 alloy after T5.

Figure 6. Basal plates in rapidly solidified Mg-1Ca isochronally annealed to 140 °C. (0002)Mg reflection.

Also the addition of Ca to some Mg alloys can lead to the precipitation of metastable or stable compounds in the form of plates parallel to the basal plane of α-Mg matrix. Very

47 thin plates of hexagonal phase were observed in Mg-1Ca-1Zn alloy both as cast and annealed at 200 °C for 24 h. after solution treatment [25]. These plates had the same structure as hexagonal basal plates in [7] and [24]. Basal plates with similar structure and size (diameter ~ 20 nm, thickness ~3 nm) were identified in Mg-2Zn-1Ca after creep at 150 °C and 64 MPa and in Mg-2Zn-0.8Ca-2RE(misch metal)-0.5Zr alloy in as cast and T6 state [26]. Basal plates of a transition phase were also observed in rapidly solidified Mg-1Ca alloy after isochronal annealing up to 140 °C [27] – see Fig. 6. These precipitates, unfortunately, dissolve at temperatures above 200 °C and so Ca cannot be used to design creep resistant Mg alloys for high temperatures application. Not only some of Mn – Re and Mn-Sc compounds precipitated in the form of plates parallel to the basal plane of α- Mg matrix, which are stable at elevated and high temperatures, can inhibit cross slip of basal dislocations and non-basal slip at those temperatures. The splitting of a dislocations into partials due to the low stacking fault energy is effective at the temperatures where these mechanisms are controlling ones. A relatively low Y and Zn contents in solid solution decreases stacking fault energy in α-Mg matrix considerably – see in Fig. 7 a contrast from stacking faults and partial dislocations accommodating thermal stresses arising during cooling down Figure 7. Stacking faults and partial basal dislocations squeeze cast MgY4Zn1Mn1 alloy. Splitting near grain boundary eutectic in as cast MgY4Zn1Mn1 of basal dislocations into partials was alloy. recently reported in Mg-0.9 mol.% Y0.04 mol.% Zn alloy [28]. The microslip bands typical for metals with low stacking fault energy were observed there after creep to the minimum creep rate at 650K and 20 MPa. Excellent creep resistance of MgY4Zn1Mn1 alloy after T5 treatment was observed at 350 °C and 30 MPa – see Table 1. Such pronounced decrease of stacking fault energy of α-Mg matrix has not yet been reported for any other alloys with combination of Zn and other REs or for those with high Zn concentration.

References [1] G.W. Lorimer in Proc. London Conf. Magnesium Technology (Ed.: H. Baker), Inst. of Metals, London, 1896, p. 47. [2] M. Ahmed, G.W. Lorimer, P. Lyon, and R. Pilkington in Magnesium Alloys and Their Applications (Eds.: B.L. Mordike and F. Hehman), DGM Informationsgesellschaft, Oberusel, 1992, p. 301. [3] B. Smola, I. Stulíková, J. Pelcová, F. von Buch, and B.L. Mordike, phys. stat. sol. (a) 2002, 191, 305 – 316.

48 [4] B. Smola, I. Stulíková, F. von Buch, B.L. Mordike, Mat. Sci. Eng. A 2002, 324, 113– 117. [5] I. Stulíková, B. Smola, and B.L. Mordike, phys. stat. sol. (a) 2002, 190, R5 –R7. [6] I. Stulíková, B. Smola, F. von Buch, and B.L. Mordike, Mat.-wiss. u. Werkstofftech. 2003, 34, 102 – 108. [7] B. Smola, I. Stulíková, J. Pelcová, F. von Buch, and B.L. Mordike, Z. Metallkde. 2003, 94, 553 – 558. [8] I.J. Polmear, Mater. Sci. Tech. 1994, 10, 1 – 16. [9] S. Kamado, Y. Kojima, R. Ninomyia, and K. Kubota in Proc. 3rd Int. Magnesium Conf. (Ed.: G.W. Lorimer), The Inst. of Materials, Cambridge, 1997, p. 327. [10] P. Vostrý, I. Stulíková, B. Smola, F. von Buch, and B.L. Mordike, phys. stat. sol. (a) 1999, 175, 491 – 500. [11] J.F. Nie, Scripta Mater. 2003,48, 1009 – 1015. [12] P.J. Apps, H. Karimzadeh, J.F. King, G.W. Lorimer, ibid, 1023 – 1028. [13] J.F. Nie, B.C. Muddle, Scripta Mater. 1999, 40, 1089 - 1094 [14] J.F. Nie, B.C. Muddle, Acta Mater. 2000, 48, 1691 – 1703. rd [15] M. Hisa, J.C. Barry, and G.L. Dunlop in Proc. 3 Int. Magnesium Conf. (Ed.: G.W. Lorimer), The Inst. of Materials, Cambridge, 1997, p. 369. [16] L.Y. Wei, G.L. Dunlop, H. Westengen, J. Mater. Sci.1996, 31, 387 – 397. [17] K.I. Gradwell, PhD thesis, University of Manchester, 1972. [18] K. Kiehn, B. Smola, P. Vostrý, I. Stulíková, and K.U. Kainer, phys. stat. sol. (a) 1997, 164, 709 – 723. [19] B.L. Mordike, F. von Buch, in Magnesium Alloys and Their Applications (Ed.: K.U. Kainer), Wiley-Vch Verlag, Weinheim, 2000, p. 35. [20] F. von Buch, J. Lietzau, B.L. Mordike, A. Pisch, R. Schmid-Fetzer, Mater. Sci. Eng. A 1999, 263, 1 – 7. nd [21] T.B. Massalski, Binary Alloys Phase Diagrams, 2 ed., American Society for Materials, Materials Park, OH, 1990. [22] B. Smola, I. Stulíková, J. Pelcová, F. von Buch, and B.L. Mordike, in Magnesium Alloys and Their Applications (Ed.: K.U. Kainer), Wiley-Vch Verlag, Weinheim, 2000, p. 92. [23] I. Stulíková, B. Smola, J. Pelcová, B.L. Mordike, “Ageing Characteristics and Creep Resistance of Mg-Y-Nd-Sc-Mn Alloy”, this conference. [24] D. H. Ping, K. Hono, F.F. Nie, Scripta Mater. 2003, 48, 1017 – 1022. [25] J.F. Nie, B.C. Muddle, Scripta Mater. 1997, 37, 1475 – 1481. [26] T. Horie, H. Iwahori, Y. Seno, and Y. Awano, in Magnesium Technology 2000 (Eds.: H.I. Kaplan, J.N. Hryn, and B.B. Clow), TMS, Warrendale, 2000, p. 261. [27] P. Vostrý, I. Stulíková, B. Smola, W. Riehemann, B.L. Mordike, Mater. Sci. Eng. A 1991, 137, 87-92. [28] M. Suzuki, T. Kimura, J. Koike, K. Maruyama, Scripta Mater. 2003, 48, 997 – 1002.

Acknowledgement The supports by German Research Foundation (DFG), by the Czech Grant Agency (GAČR project 106/03/0903) and by Ministry of Education of Czech Republic within the framework of the research program MSM 113200002 are gratefully acknowledged.

49

Regularities in the Phase Diagrams and Behaviour During Aging of the Ternary Mg-Base Alloys Containing Two Rare-Earth Metals of Different Subgroups Lazar Rokhlin, Tatyana Dobatkina, Nadezhda Nikitina (Baikov Institute of Metallurgy and Materials Science, Moscow)

The rare earth metals (RE) are well-known as effective alloying elements for magnesium alloys. A specific feature of the rare earth metals as they are added to magnesium is significant difference between them. Thus, neodymium shows higher strengthening effect than cerium and yttrium shows higher strengthening effect than neodymium. As it was established the best mechanical characteristics of magnesium alloys could be achieved when two separated rare earth metals are added to magnesium with each of them belonging to different subgroups, the cerium or yttrium ones. One of this kind systems is MgNd-Y. A cast commercial alloy WE54A on base of it shows the highest strength at elevated temperatures. On base of this system a number of the Russian commercial magnesium alloys with high strength properties were developed, as well. For example, the cast alloy ML19 [1, 2]. The strength properties of Mg alloys with rare earth metals are connected at a great extent with the Mg-RE phase diagrams, especially, solubility of the rare earth metals in solid Mg and ability of Mg supersaturated solid solution to decompose during aging. The regularities in the ternary phase diagrams of Mg with two rare-earth metals belonging to different subgroups and their behavior during solid solution decomposition may be revealed, as well. They base on results of the investigations on the systems Mg-Nd-Y, MgLa-Y, Mg-Ce-Y, Mg-Sm-Y, Mg-Sm-Gd, Mg-Sm-Er alloys conducted in Baikov Institute of Metallurgy and Materials Science (Moscow, Russia. Mg-rich parts of all studied ternary phase diagrams with two rare-earth metals of different subgroups show similar features. They can be seen in Figure 1 where a partial section of the Mg-Sm-Er phase diagram is shown, as an example [3]. The main features of the phase diagram are. 1) The total solubility of both rare-earth metals in solid Mg is intermediate between those of the rare-earth metals in the binary systems. It can be concluded from the boundary of the Mg solid solution field which is approximately straight line connecting points of solubility in the binary systems. 2) Only two phases are in equilibrium in solid Mg. Both of them are the compounds of the respective binary systems. 3) Each of the binary compounds can dissolve significant contents of the other rare earth metal. This fact is demonstrated by the deviation of the straight lines dividing twophase and three-phase fields from the concentration axes. The same main features are inherent to all other ternary phase diagrams of this kind studied. In order to support this viewpoint in Figure 2 Mg part of other ternary phase diagram, Mg-Sm-Y system is reproduced [4]. It show the same form of Mg solid solution area, existence only two phases in equilibrium with it and extended solid solutions on the

50 4 (Mg)+Mg Sm 5 + 5 41 Mg24Er 5

At

.%

Sm

(Mg)+ Mg41Sm 5

2 (Mg)+ Mg24Er5

0

2

(Mg)

4 At. % Er

6

8

Figure 1. Partial isothermal section of the Mg-Sm-Er phase diagram at 500 °C.

20 Mg24Y5 15

At. %

Y

b 10 e1

c

5 a1 Mg 0

a2

U a 5

e2

10 15 At. % Sm Mg Sm 5 41

20

Figure 2. Mg-Sm-Er phase diagram. Projection of the four-phasde equilibrium plane and the connecting double saturation lines

bases of the Mg-richest compounds in each of the adjoining binary systems determined by local spectral analysis. Investigations of the binary systems showed that kinetics and phase transformations during solid solution decomposition in Mg alloys with the rare earth metals are similar if the rare earth metals belong to the same subgroup (cerium or yttrium one) and different if the rare earth metals belong to the different subgroup. It is reasonable to remind the main features of them [5]. The remarkable feature of the solid solution decomposition in the alloys with elements of the cerium subgroup is the steep increase of hardness already after short aging time. Then hardness maximum is reached. In contrast with this in the alloys with elements of yttrium subgroup the hardness increase during solid solution decomposition takes place

51 only after some latent period. Phase transformations in the alloys with elements of the cerium subgroup are characterized mainly by formation of the GP-zones at initial stage and ordering of the Mg3Cd type then. This structure corresponds to the hardness maximum. At higher temperatures the plate-like precipitates are formed on the prism planes of the (100) type with three-fold symmetry. Phase transformations in the alloys with elements of the yttrium subgroup are characterized by ordering of the same Mg3Cd type at latent period and at hardness maximum the precipitates with orthorohmbic lattice are formed. They are plate-like particle disposed with three-fold symmetry on the (110) planes of the lattice. The equilibrium phase is formed as coarse plate-like particles on the (100) prism planes. The phase diagrams enable to suppose that during solid solution decomposition in the ternary alloys two successions of the phase transformations run simultaneously with each of them ending by the respective equilibrium compound. In according with this, the properties change during solid solution decomposition in the ternary system is supposed to be as if superposition of properties change resulting from both transformation successions in structure. Nevertheless, it not the case indeed in the ternary systems and in this sense the behaviour of the ternary alloys during solid solution decomposition is unusual. Unusual behaviour of the ternary alloys during solid solution decomposition was observed earlier in the Mg-Nd-Y alloys of the WE54 type [6] In this paper, additional data confirming this fact for other concentrations in the same system and in other systems are presented. In Figure 3 the representative hardness curves for the solid solution decomposition are presented for the alloys of the Mg-Nd-Y system. Ageing temperature is 250 °C. Two concentration of Y is used – 4.5 and 8.5 %. The alloys were solution treated by quenching from 510 °С / 6 h into cold water and then aged isothermally at 250 °C. The curves show certainly the hardness maximum in the ternary alloys is higher and at ageing times to be longer than that in the binary Mg-Nd alloy with the same Nd contents. The hardness increase in the binary Mg-Y alloys with the same Y contents at 250 °C is insignificant. 120

Hardness, HB (kg.mm -2)

Mg-3%Nd-8.5%Y

Mg-3%Nd-4.5%Y 80 Mg-8.5%Y Mg-3%Nd

Mg-4.5%Y 40 Quenched 1

10 Ageing time, hours

Figure 3. Hardness of the Mg-Nd-Y alloys versus aging time at 250 °C.

100

52

Resistivity, mcOm.cm

30

Mg-8.5%Y 20 Mg-3%Nd-8.5%Y 10 Mg-3%Nd 0 Quenched

1

10

100

Aging time, hours Figure 4. Electrical resistivity of the Mg-Nd-Y alloys versus aging time at 250 °C.

130 Hardness, HB (kg.mm -2)

Mg-5.2%Sm-5.2%Y

200 oC

110 o

220 C 90 200 oC 70

o

Mg-4.3%Sm-0.6%Y

220 C

50 Quenched 1

10 Aging time, hours

100

Figure 5. Hardness of the Mg-Sm-Y alloys versus aging time.

Results of the hardness measurements at ageing temperature 250 °C are supported by the electrical resistivity measurements at the same ageing temperature presented in the next Figure 4. They show as resistivity decrease significantly more depletion of Mg solid solution during ageing in the ternary alloys as compared with the respective binary alloys. This effect especially obvious in the ternary alloy with 8.5% Y. Results for the alloys with 4.5%Y are missed in the Figure 4 to make them more evident. In general the results for the Mg-Nd-Y alloys presented agree with those obtained in [6].

53 Figure 5 demonstrates the hardness curves for the ageing in the alloys of the third system Mg-Sm-Y. The alloys were preliminary hot extruded and solution treated by quenching from 520 °C / 2 h into cold water. The curves confirm the general tendency - increase of hardness maximum and shift of it to longer ageing time with increasing contents of the element of the yttrium subgroup. Investigations of the Mg-Sm-Y alloys showed a possibility to improve significantly tensile strength of the binary Mg-Sm ones by addition of Y. The results of the tensile tests of the hot extruded alloys after aging at 200 °C for 12 h immediately after hot extrusion (T5) and after intermediate solution treatment (T6) are shown in Table 1. Table 1. Tensile properties of the hot extruded Mg-Sm-Y alloys.

Temper

T5 T6

Test temperature, °C 20 250 20 250

Mg-7%Sm TS, MPa TYS, MPa 263 212 185 152 274 159 220 132

Mg-7%Sm-3%Y E, % 13 26 13 20

TS, MPa TYS, MPa 340 308 272 221 329 249 259 195

E, % 6 21 9 10

Studying structure of the ternary alloys after solid solution decomposition we did not observe any new precipitates except those had been observed earlier in the binary systems. In accordance with the phase diagrams both type precipitates which are inherent to the alloys with the rare-earth metals of different subgroups were existed in structure together. At aging times near hardness maximum the coarse precipitates were identified as those of the rare-earth metals of Ce subgroup and fine precipitates were identified as those of the rare-earth metals of Y subgroup. Around the coarse particles the precipitate free zones (PFZ) of the fine particles were watched [7]. Taking into consideration nature of the precipitate free zones their existence may be considered as a proof of dissolution of the rare-earth metal of Y subgroup in the precipitates being inherent to the alloys with the element of Ce subgroup.

Conclusions 1. Ternary phase diagrams of Mg with two rare-earth metals of different subgroups show similar features which in the Mg part consist of total solubility of both rare-earth metals being intermediate between those in the binary systems, existence in equilibrium with solid Mg only two Mg-richest compounds in Mg-RE binary systems and an ability of each of the rare-earth compounds to dissolve significantly other rare earth metal. 2. Kinetics of Mg solid solution decomposition and accompanying strengthening effect can be significantly improved by using as alloying elements two rare earth metals of different subgroups (cerium and yttrium ones). In the ternary alloys of this kind better strength properties can be reached.

54 3. Change of the Mg solid solution decomposition in the ternary alloys as compared with those in the respective binary ones can be explained by solubility of every rare earth metal in the precipitates of the other formed during the process.

References [1] Standard Specification for Magnesium Alloys in Ingot Form for Sand Castings, Permanent Mold Castings, and Die Castings. American Society for Testing and Materials. Designation: B93/B93M - 91. [2] Casting magnesium alloys. Grades. USRR State Standard No 2856-79 (in Russian). [3] L.L. Rokhlin, N.I. Nikitina, Metally, 2002, (2) 119-123 (in Russian). [4] M.E. Drits, L.L. Rokhlin, N.P. Sirchenko, Izv. Akad. Nauk SSSR, Met., 1983 (6), 78-82 (in Russian). [5] L.L. Rokhlin. Magnesium Alloys Containing rare Earth Metals. Taylor and Francis. London – New York, 2002. [6] P. Vostry, T. Stulikova, B. Smola, M. Cieslar, B.L. Mordike, Z. Metallkd., 1988, 79(5), 340-344. [7] N.P. Abrukina, in “Magnivye Splavy dlya Sovremennoy Tekhniki” (Magnesium Alloys for Modern Technics), “Nauka”, Moscow (Ed. N.P. Lykishev), 1992, 82-86 (in Russian).

55

Development of New Magnesium Alloys for Advanced Applications B. Bronfin1, E. Aghion1, F. von Buch2, S. Schumann2, H. Friedrich2 1 2

Dead Sea Magnesium Ltd, Beer Sheva, Israel Volkswagen Ag, Wolfsburg, Germany

1

Introduction

To date a significant interest is being focused on the development of new die casting alloys for high temperature applications, especially for the automotive industry. This interest relates mainly to powertrain components such as gearbox housing, oil pan, transfer case, crankcase, oil pump housing, intake manifold and others. These parts operate at various elevated temperatures and therefore enhanced creep resistance and bolt load retention properties are essential for adequate performance. Unfortunately, the common alloys of AZ and AM series are not suitable for manufacturing drive train parts operating at temperatures higher than 130 °C. The insufficient creep resistance of regular commercial alloys used for manufacturing of various housings can result in as clumping load reduction of bolted joints that causes poor bearing-housing contact. This leads to oil leaks and increased noise and vibration. In addition, the creep resistant alloys developed in the 70’s (AS21 and AS41) and in the late 80’s (AE42) also did not find their place in powertrain applications due to various reasons such as poor castability and corrosion resistance, increased cost and low strength [1]. Most of the drive train components for regular cars are being produced by high-pressure die casting – a high volume production technology. In this technology, the alloy cost represents a significant portion of the component cost. Therefore the alloy cost becomes a major factor in the alloy development, which imposes on new alloys the importatnt requirement to be price competitive to existing magnesium and aluminum commercial alloys.

2

Metallurgical background for the development of high temperature diecasting magnesium alloys

The requirements of the affordable cost and good die castability of a new alloy reduce possible options to alloying systems, which contain Al or Zn as major alloying elements and Mn, Si, Ca, Sr and Ce-based mishmetal as relatively small additions. When developing Mg alloys for die casting applications it should be taken into consideration that alloying with Al is strongly recommended in order to provide good fluidity properties (castability). Therefore, a magnesium alloy should countain a sufficient amount of Al in the liquid state prior to solidification. However, the presence of Al results in the formation of eutectic Mg17Al12 intermetallics, which adversely affect the creep resistance.

56 Hence, additional alloying elements that can form specific intermetallics with alluminum should be added into an Mg-Al alloy in order to suppress the formation of β-phase (Mg17Al12). Considering a hypothetical ternary phase diagram for the Mg-Al-Me system (Me- is the third alloying element) let’s assume that in this system three intermetallic compounds can generally form: Mg17Al12, MgxMey, AlzMew (Fig. 1). In order to suppress the eutectic reaction involving the formation of the β-phase (Mg17Al12) the element Me should react with aluminum to form the intermetallic compound AlzMew. In this case the pseudobinary section Mg-AlzMew will be active. This will take place only in the case when the affinity of Me to Al is higher than that of Mg and the formation of AlzMew is preferential to the formation of the MgxMey intermetallic compound. Analysis of available binary phase diagrams Mg-Me and Al-Me [2] have shown that only rare earth elements, alkaline earth elements and 3-d transitions elements can comply with the above requirements. Another option is to reduce the Al content to relatively low level in order to prevent the formation of β-phase and in parallel introduce elements with high affinity to Mg to provide formation of stable phases that can impede grain sliding. This concept was used in the development of alloys such as AS21 and AS41 [1], where formation of Mg2Si intermetallics improves creep resistance compared to Az and Am alloys.

Figure 1. Hypothetic ternary phase diagram Mg-Al-Me

This approach was further expanded by Hydro Magnesium in the delopment of the alloy named AS21X [3] and by Magnesium Electron in the development of MEZ alloy [4]. Alternative attempts in the development of new creep resistant alloy are based on the idea of suppressing the formation of β-phase in alloys containing 5–9% Al. This was obtained by introducing alkaline earth and rare earth elements to bind excess of aluminum into the stable intermetallic compounds. As an outcome of these attempts several new Mg alloys for high temperature applications were recently developed with the intention of obtaining an adequate combination of different properties such as creep resistance, castability, mechanical properties, corrosion resistance and affordable cost [5–8]. Unfortunately, this combination of properties is difficult to achieve and evidently most of the new alloys can only partially address the required properties and performances. Most of the above reviewed die cast Mg alloys contain various concentrations of Al, Ca, Sr, and rare earth (RE) elemts. Alkaline earth elements and Rare earth were used to improve the creep resistance while Al was kept to provide adequate castability.

57 The final cost of the new alloy depends on the concentration of the alloying elements and their price. In fast, the use of different concentrations of the above alloying elements resulted in positive modifications of some properties while other properties were deteriorated. This was also followed by a slight increase in alloy costs. A schematic illustration showing the effect of Al and alloying elements on castability, properties and cost of magnesium based die casting alloys is given in Fig. 2a. This diagram illustrates that increasing Al content in magnesium alloys results in better castability and increased ambient tensile, compressive and fatigue strength as well as improved corrosion resistance. On the other hand, creep properties, ductility and impact strength are significantly deteriorated (Fig. 2a). Alloying with Al has little affect on alloy cost. In constrat to this, alloying with alkaline earth elements (Ca, Sr) and rare earth mishmetal (RE) leads to deterioration of castability, ductility and impact strength as well as increase in cost. The benefit of this alloying is mainly associated with significant improvement of creep resistance followed by moderate increase of ambient strength (Fig. 2b).

Figure 2. Effect of Al (a) and alloying elements Ca, Sr, RE (b) on properties and cost of magnesium die casting alloys

Calcium has the atomic weight, which is half of that of atomic weight of Sr and RE elements. Therefore, significanty less Ca additions are required compared to Sr and RE additions in order to provide the same volume percentage of the second phase (Table 1). Table 1. Comparative effect of Calcium, Strontium and Rare Earth Metals on the formation of Intermetallic compounds

Element

Ca

Sr

Al2Ca

Al2Sr

Al4Sr

Al11RE3

Al2RE

2.35

2.98

2.85

4.02

4.95

Relative quantity of the alloying 1 element to form equal weight percentage of precipitates

2.2

1.1

1.9

3.5

Relative quantity of the alloying 1 element to form equal volume percentage of precipitates

2.6

1.4

3.25

7.4

Possible precipitate 3

Density [g/cm ]

RE

58 Consequently, alloying with Ca provides maximum increase of both ambient tensile and compressive yield strength and especially creep strength for the equal weight percentage additions compared to Sr and RE mishmetal. Since Ca is a relatively inexpensive element, it also provides the maximum increase of the above properties in terms of value to cost. However, it is also evident that alloying with Ca in concentration more than 0.3% causes castability and ductility deterioration. Hence, the actual challenge is to select optimum combination of alloying ingredients (Al, Ca, Sr, RE) to provide the best service properties combined with affordable cost.

3

DSM-VW development of creep resistant alloys

In order to address the required combination fo properties, DSM and Volkswagen AG initiated a comprehensive alloy development program aiming at developing new creep resistant and cost effective diecasting magnesium alloys. As an outcome of this work two new Mg alloys designated MRI 153M 230D, have been developed [9–11]. MRI 153M is a Beryllium free, low cost creep resistant alloy with the capability of long-term operation at temperatures up to 150 °C under high stresses. MRI 153M alloy exhibits die castability similar to that of AZ91D alloy. This was demonstrated and confirmed by die casting trials of complicated components such as transmission housings for VW vehicles: Polo, Lupo, Golf, Audi A3 ad Audi A6 [9–12]. Various properties of MRI 153M alloy in comparison with commercial alloys AZ91D, AE42 and AS21 are summarized in Table 2 and Figs. 3–4. It is evident that MRI 153M has corrosion resistance and mechanical properties at room and elevated temperatures similar or better than AZ91D and other commercial alloys. Table 2. Mechanical and corrosion properties of diecasting alloys MRI 153 M and MRI 230D

Properties

MRI 153M MRI 230D AZ 91D AE42

AS21

Tensile Yield Strength (MPa)

At 20 °C 170 At 150 °C 135

180 150

160 105

135 100

125 87

Ultimate Tensile Strength (MPa)

At 20 °C 250 At 150 °C 190

235 205

260 160

240 160

230 120

Elongation (%)

At 20 °C 6 At 150 °C 17

5 16

6 18

12 22

16 27

Compression Yield Strength At 20 °C 170 (MPa) At 150 °C 135

180 150

160 105

115 85

110 80

Impact strength (J)

8

6

8

12

14

120

110

100

80

75

0.09

0.10

0.11

0.12

0.34

(1)

Fatigue strength (MPa) Corrosion Rate (mg/cm2/day) (1)

(2)

Rotating bending 5 ×10 Cycles 7

(2)

200 hr. Salt Spray Test (ASTM Standard B-117)

59 The creep properties of MRI 153M, at 130–150 °C under stresses of 50–85 MPa, are also significantly superior to those of commercial alloys. This may introduce MRI 153M as a superior choice for automotive applications such as transmission housing, oils pans, valve covers, intake manifolds, etc.

Figures 3. Creep resistance of MRI 153M compared to conventional alloys at 135 °C under stress of 85 MPa

Figure 4. Creep resistance of MRI 153M at 150 °C in Comparison with Conventional Alloys

MRI 230D is a diecasting alloy, which was developed to address automotive applications such as engine blocks operating at higher temperatures up to 190 °C. The alloy has excellent creep resistance combined with good castability, high strength and superior corrosion behavior. The typical properties of MRI 230D alloy are summarized in Table 2 and in Fig. 5 in comparison with commercial magnesium alloys and aluminum alloy A380. The results obtained show that MRI 230D and A380 exhibit similar tensile creep behavior at 150– 175 °C under stress of 70 MPa.

60

Figure 5. Creep Resistance of MRI 230D in Comparison with Conventional Alloys at 175 °C/70 MPa

Table 3 illustrates the effect of different aging conditions on tensile properties and minimum creep rate (MCR) of both MRI’s alloys. It is evident that direct aging provides a wide combination of strength and ductility without deterioration of creep properties. Hence, it can be considered as a powerful tool for optimization of the required properties. Table 3. Effect of direct aging (T5) on tensile properties and minimum creep rate (MCR) of MRI 153M and MRI 230D alloys.

Alloy

MRI 153M MRI 230D

State

F T51 T52 T53 F T51 T52 T53

TYS

UTS

MPa

MPa

170 177 168 166 180 195 181 166

250 240 255 258 235 240 245 270

E%

6 4 7 8 6 4 5 8

MCR ⋅ 10–9, s–1 150 °C 70 MPa

175 °C 70 MPa

2.51 2.63 2.47 2.62 – – – –

– – – – 0.95 0.93 0.92 0.96

These findings enable to extend the range of applications for MRI’s alloys and validate their flexibility and feasibility to tailor properties according to operating requirements for different components.

4

Conclusions

Metallurgical principles underlying selection of new pratical diecasting alloys for high temperature applications were considered and discussed. The proposed approach was implemented at the development of new creep-resistant diecasting alloys MRI 153M and MRI 230D. Direct aging can be applied to optimize properties and performances of new alloys.

61

5

References

[1] J.F. King, Magnesium alloys and their applications (Eds.: B.L. Mordike and K.U. Kainer), Weinheim, 1998, 37-47. [2] Baker et al., Metals Handbook, vol. 2, Metals park, OH: American Society for Metals, 1992. [3] K. Pettersen etal., Magnesium Alloys and Their Applications, (Ed.: K.U. Kainer) Weinheim, 2000, 29-34. [4] P. Lyon, J.F. King, and K. Nuttal, Proceedings of the Third International Magnesium Conference, (ED.: G. W. Lorimer) Manchester, 1996, 99-108. [5] S. Koike et al., SAE Technical Paper 2000-01-1117, Detroit, MI, 2000. [6] B. Powell et al., SAE Technical Paper 2001-01-0422, Detroit, MI, 2001. [7] B. R. Powell et al., Magnesium Technology 2002, (Ed.: H.I. Kaplan) Warendale, PA: TMS, 2002, 123-130. [8] M. Pekguleryuz and E. Baril, Magnesium Technology 2001, (Ed.: J.N. Hryn), Warrendale, PA: TMS, 2001, 119-126. [9] B. Bronfin et al., Magnesium Technology 2001, (Ed.: J.N. Hryn) Warrendale, PA: TMS, 2001, 127-130. [10] F. von Buch et al., “New Die Casting Alloy MRI 153 for Power Train Applications”, Magnesium Technology 2002, ed. H.I. Kaplan (Warrendale, PA: TMS, 2002), 61-68. [11] E. Aghion et al., SAE Technical Paper 2003-01-191, Detroit, MI, 2003. [12] H. Friedrich and S. Schumann, Procceedings of the Second Israeli International Conference on Magnesium Science & Technology, (Ed.: E. Aghion and D. Eliezer), Dead Sea, Israel, 2000, 9-18.

62

Creep Resistance and Creep Kinetics of Mg-alloys W. Blum1, Y. J. Li1, X. H. Zeng1, B. von Großmann2, C. Haberling2, H.-G. Haldenwanger2 1 2

Universität Erlangen-Nürnberg, Inst. f. Werkstoffwissenschaften LS1, 91058 Erlangen, Germany Audi AG, 85045 Ingolstadt, Germany

1

Introduction

Resistance against creep (inelastic deformation ε with time t under stress σ at absolute temperature T) is an important issue in application of Mg-base light-weight alloys at temperatures up to 200 °C. However, the data base on creep properties of Mg-alloys is rather poor. Only few ε-t curves for constant σ and T have been published. These are often limited to primary creep at extremely low rates in a small range of stresses. This will be demonstrated with examples from the literature. In fact, creep should be treated as an integral part of inelastic deformation, to be quantified by the relation between σ, ε (inelastic) strain rate ε& , and T and furthermore to be linked to microstructure an its evolution. This concept has been applied before to several conventional Mg-base alloys [1, 2]. In this paper we extend it to the new creep resistant alloy AJ52. We will demonstrate that its creep resistance reaches or even tops that of most other Mg-base alloys.

2

Visualization of creep

2.1

Strain-time curves

Limitation of testing time naturally leads to curves like those displayed in Fig. 1 taken from [3]. Curves at high stress are steep and extend to large strains. Curves for low stress, on the other hand, have a small slope and extend to smaller strains than the former ones. Only those curves which happen to lie at an angle between 30 an 60° are easily evaluated for the slope ε& . This makes the linear ε-t diagramm unsuitable to fisplay the results for a large range of creep conditions (σ, T). 2.2

Creep rate-strain curves

Plotting creep data as log ε& versus strain ε is a means to arrive at curves with similar shape [4, 5] which are easy to compare over a large range of σ and T and clearly show whether or not the relative minimum in ε& has been passed. The creep rate is obtained from the ε-tcurves by determining ε& = Δε/Δt where Δε must exceed the experimental error of ε determination which is typically in the order of 10-5 to 10-4. Curves form the literature haven been evaluated by averaging the measured ε-t data until a sequence of εi-ti was obtained so that Δε i = ε i +1 − ε i −1 > 10 −4 . The symbols in Fig. 1 show the ε-t data pairs resulting from this procedure. The creep rate ε&i (ε i ) was determined as ( ε i +1 − ε i −1 ) / ( ti +1 − ti −1 ) . Fig. 2 displays the curves so obtained from the curves of Fig. 1. The decrease of the ex-

63

Figure 1. Creep strain ε as function of time t for AZ91, AS41 and ZACE05411 (Mg-0.5Zn-4Al-1Ca1RE) at 175 °C; from [3].

Figure 2. Creep rate strain ε .

ε& evaluated from Fig. 1 versus

plored strain interval with decreasing creep rate resulting from testing time limitation is evident. we note in passing that it is easily and accurately possible to retur to the ε-t curves by integrating the ε& -ε- curves.

3

Review of recent data from literature

A number of published creep curves have been evaluted as desribed in the previous section. Comparison of creep data obtained at different T is facilitated by multiplying ε& with exp (Q/(RT)), where R is the gas constant and Q is an appropriate value of the effective activation energy of creep, to arrive at the so-called Zener-Hollomon parameter Z = ε& ⋅ exp (Q /( RT )) . The transition from ε& to Z means a vertical shift of the log ε- ε& curves. If Q is chosen appropriately, the log Z-ε curves depend mainly on σ, and are virtually independent of T within a limited T-range. This has been demonstrated for serval MgAl-base alloys in [1, 2] and is also apparent from Fig. 4 below. The choice of Q is rather uncritical. In agreement with [2] Q was set to 135 kJ/mol. Fig. 3 displays the minimum values Zmin of Z derived from the published curves. It is seen that most of them are grouping in between the data fields for AZ91 and AS21 reported in [2]. Most of the symbols are marked by an arrow indicating that ε& had not yet reached its relative minimum which means that the were terminated in the primary stage of creep. Most of the data stem from tensile tests where fracture may cause an extra increase in creep rate [13]. It is also evident that most data habe ben acquired in a

64

Figure 3. Zmin as function of σ for AZ91 (data band from [2], symbols for 125 – 175 °C [3, 6, 7, 8]), AZ51 (200 °C [9]), AS21 (data band from [2], symbols for 125 – 175 °C [6, 7, 8]), AS41 (175 °C[3]), AM50 + Ca (150 °C [10]), AE42 (150 – 175 °C [3, 6, 7]), MEZ (150 °C [11]), MRI 153M (135 – 150 °C [7]), MRI 230D (150 °C [7]), ZA85 (170 °C, curves available in [12]), ZACE 05613 (Mg-0.5Zn-6Al-1Ca-3RE, 175 °C [3]), ZACE 05411 (175 °C [3]), ZE41 (200 °C [7]). Arrows indicate that creep test ended before a relative minimum was reached.

relatively narrow window (corresponding to low minimun creep rates < 10-6/s). Thus the data base obtained from literature appears to be rather weak and insufficient as a basis for modelling the creep behavior in the full range of loading conditions.

4

Creep behavior of AJ52

AJ52 is a new temperture creep resistant Mg-Al-Sr alloy (composition/mass%: 4.5– 5.7 Al, 1.7–2.3 SR, 0.3–0.6 Mn). It was produced by die casting a plate with stepped thickness profile. Specimens were prepared from the portion of the plate with 6 mm thickness. Testing was done in uniaxial compression at constant mechanical (elastic plus inelastic) strain rate ε&mech as well as at constant nominal stress σ (load per average cross section at test temperature). The relation between σ , ε& and ε is displayed in Fig. 4. It is seen from

65

−3

Figure 4. Relation between ε& , σ and ε for AJ52. Tests 1 to 4 at constant rate ε&mech = 10 /s (thin horizontal lines in lower part) yield the σ − ε curves shown as thick lines in the upper part. Creep tests a to h where σ was held constant after loading (thin horizontal lines in upper part) yield the Z-ε curves shown as thick lines in the lower part. Symbols in upper part show ε (σ) measured near the end of the loading phase of the creep tests when the rate had declined to 10 −3 /s

66 the upper part of that the yield stress at ε&mech = 10 −3 / s is about 100 MPa in agreement with [14]. The symbols in the upper part of Fig. 4 mark the strains ε obtained by the end of loading the creep specimens at rates 10 −3 / s < ε&mech < 4 ⋅ 10 −2 / s when ε& had fallen to 10−3 / s. It is seen that the symbols fall close to the σ-ε curves for at ε&mech = 10 −3 / s . This confirms the result reported previously [1, 2] that the strain to reach a certain deformation resistance ε&, σ is not very sensitive to the preceding ε&mech -history. Work hardening is seen in the σ −ε& curves at constant rate as increase in flow stress σ and in the ε&−ε curves at constant stress as decrease in rate. The maximum deformation resistance is characterized by the maximum flow stress at constant rate and the minimum Z at constant stress. Fig. 5 shows the result. The symbols for the maximum deformation resistance fall below the range of data shown in Fig. 3. This means that the creep resistance of AJ52 is quire high in relation to that of the alloy shown in Fig. 3. In particular, it is superior to that of AS21. The relatively high creep resistance of AS21 has been explained by particle hardening due to Mg2Si precipitates [15]. A similar explanation holds for AJ52 where Sr-containing intermetallic precipitates have been observed [16], Thus AJ52 is distincly superior in creep resistance to AZ91 with only slight disadvantage in yield stress at temperatures below 150 °C [14].

Figure 5. Maximum deformation resistance (minimum Z, maximum σ) of AJ52 in comparison to the field of data from Fig. 3.

67

5

Conclusions

1. Creep rate-strain curves are needed for full description of creep behavior. 2. Temperature compensation of creep rate by an arrhenius term eliminates their temperature dependence in a limited range of temperatures indicating that the structural development in the alloy is mainly dependent on stress. The Sr-containing alloy AJ52 belongs to the most creep resistant Mg-Al base alloys at temperatures below 200 °C.

6

References

[1] B. Watzinger, P. Weidinger, F. Breutinger, W. Blum, R. Rösche, H. Lipowsky, and H.-G. Haldenwanger in Magnesium Alloys and their Applications (Ed. B. L. Mordike and K. U. Kainer), MATINFO Werkstoff-Informationsgesellschaft, Frankfurt, 1998, 259–264. [2] W. Blum, P. Zhang, B. Watzinger, B. v. Grossmann, and H.-G. Haldenwanger, Mater. Sci. Eng. A 2001, 319–321, 735–740. [3] S. Kamado, J. Koike, K. Kondon, and Y. Kawamura, Mater. Sci. Forum, 2003, 419– 422, 21–34. [4] W. Blum and B. Ilschner, phys. stat. sol., 1967, 20, 629-642. [5] W. Blum, in materials Science and Technology (Ed. R. W. Cahn, P. Haasen an E. J. Kramer), Vol. 6, Plastic Deformation and Fracture of Materials (Ed. H. Mughrabi), VCH Verlagsgesellschaft, Weinheim, 1993, 359–405. [6] M. S. Dargusch, G. L. Dunlop, and K. Pettersen, in Magnesium Alloys and their Applications (Ed. B. L. Mordike and K. U. Kainer), MATINFO WerkstoffInformationsgesellschaft, Frankfurt, 1998, 277–282. [7] E. Aghion, B. Bronfin, D. Elienzer, F. von Buch, S. Schumann, and H. Friedrich, Mater. Sci. Forum, 2003, 419–422, 407–418. [8] I. A. Anyanwu, Y. Gokan, S. Nozawa, S. Komado, Y. Kojima, S. Takeda, and T. Ishida, Mater. Sci. Forum, 2003, 419–422, 445–450. [9] G. Yuan, M. Liu, W. Ding, and I. Akihisa, Mater. Sci. Forum, 2003, 419–422, 425– 432. [10] Y. Terada, N. Ishimatsu, R. Sota, and K. Ohori, Mater. Sci. Forum, 2003, 419–422, 459–464. [11] I. P. Moreno, T. K. Nandy, J. W. Jones, J. E. Allison, and T. M. Pollock, Scripta Mater., 2003, 48, 1029–1034. [12] M. Vogel, O. Kraft, and E. Arzt, Scripta Mater., 2003, 48, 985–990. [13] W. Blum, P. Weidinger, B. Watzinger, R. Sedlacek, R. Rösch, and H.-G. Haldenwanger, Z. Metallkde., 1997, 88, 636–641. [14] M. Ö. Pekgüleryüz, Technical report, 2003, Noranda Inc. Technology Centre, Canada. [15] P. Zhang, R. Agamennone, W. Blum, B. v. Grossmann, and H.-G. Haldenwanger, in Proc. Int. Symp. on Magnesium alloys and Their Applications (Ed. K. U. Kainer), Wiley-VCH, Weinheim, 2000, 716–721. [16] Technical report, 2003, http://www.noranda.com

68

Creep Behaviour of MgZnRE Alloys Formed by Semi-Solid-Casting E. D. Morales, N. Hort, K. U. Kainer GKSS Research Center Geesthacht GmbH, Geesthacht, Germany

1

Introduction

Magnesium alloys have been the target of the automotive industry in the last decades because of their low density, high specific strength, excellent machinability, high stiffness and good castability [1, 2, 3]. However, special focus has been given to applications at high temperatures, like on power trains or transmission engines, where conventional magnesium alloys are limited due to their poor strength and creep resistance. Due to the increasing demand on stable materials at high temperatures, new alloys have been developed to overcome these limitations. Quite often, these new alloys, like the MgZnRE alloys investigated in this work, present limited castability during high pressure die casting (HPDC). Another limitation for the use of the HPDC process is that in thick-walled sections porosity is found frequently because of the turbulent die filling and solidification shrinkage [1]. Therefore, other casting processes like semi-solid forming are under research for magnesium base materials. Thixocasting constitutes an alternative forming process because of the near-net-shapes formed parts, the lower energy consumption and the finer microstructures obtained. As cited in [2], the development of high temperature resistance magnesium alloys requires of alloying elements with a high maximum solubility in magnesium falling sharply with temperature and strengthening phases rich in magnesium with a high decomposition temperature. Rare earth additions on magnesium alloys fulfil these requirements; however, it has been found that Ca additions have also similar effects, being preferable over the RE additions because of the lower costs [2, 4]. Likewise, Zn added to Ca-Mg alloys increases tensile strength, ductility and creep resistance at elevated temperatures [4]. The objective of the present work is to study the effect of a semi-solid process (thixocasting) on the microstructure and creep behaviour of two Mg-Zn alloys. One of the alloys is the commercial MEZ, patented by Magnesium Elektron, and the other is an MEZ-type modified alloy with Ca-additions (ZEC). Previous studies have reported results on Mg-Zn-Ca alloys produced by rapid solidification [7] and die-casting [4], but the effect of rare earths on such alloys has not been understood. Besides, the use of semi-solid processes on aluminium free Mg-Zn alloys is until now unpopular. This work attempts to emphasize the feasibility of new forming processes on such aluminium-free systems.

2

Experimental

The chemical composition of the studied alloys was determined from the cast plates by means of the inductively coupled plasma technique (ICP). The results are shown in Table 1. Stepped

69 plates of both alloys with dimensions of 110X250 mm and thicknesses of 2, 6, 10 and 14 mm were produced by thixocasting using a commercial die-casting machine. To produce the semi-solid plates, extruded bars of both materials supplied by Magnesium Elektron were heated by induction during 8 to 9 minutes and without stirring to a temperature between liquidus and solidus and injected into a preheated mould. The semi-solid region 60 % liquidus – 40 % solidus was estimated using processing temperatures of 615 °C for the ZEC alloy and 640 °C for the MEZ alloy. To achieve a globular microstructure, the raw material was heated 25 °C above the 60 % liquidus – 40 % solidus region, cooled to the operation temperature and poured immediately into the sleeves of the casting machine. Tensile and creep specimens were machined from the 10 and 14 mm thicknesses sections. Tensile tests were performed at room temperature using specimens with 6 mm diameter and strain rates of 10-3 s-1. Three specimens were examined for every alloy. Tensile creep tests were conducted on specimens with a gauge diameter of 5 mm and a gauge length of 30 mm using constant stress creep machines at temperatures of 135, 150, 175 and 200 °C and stresses of 60, 70 and 80 MPa. The specimens crept during 200 hours maximum. Table 1. Chemical Compositions of the MgZnRE Alloys

Alloy ZEC MEZ™

Al 0.06 <0.01

Ca 3.5 –

Fe 0.009 0.001

Mn 0.01 0.2

Ni 0.002 0.002

Si 0.03 <0.01

Sr 0.1 –

Zn 3.6 0.5

Zr <0.05 0.1

RE 0.7 3.1

The examination of the microstructure of the semi-solid alloys in the as-cast state was performed with the use of a scanning electron microscope (DSM 962 Zeiss). The specimens were ground with SiO2 grinding paper and polished up to 0.05 μm SiO2 (OPS™). The chemical composition of the phases present in the alloys was determined by means of energy dispersive X-ray spectroscopy (EDX-WINEDS Version 3.0).

3

Results and Discussions of Results

3.1

Microstructure

The microstructure of the thixocast alloys consists of globular primary α-Mg grains surrounded by secondary α-Mg grains and intermetallic phases. The primary rounded α-Mg grains are formed when the raw material is heated to a temperature between liquidus and solidus and correspond to the solid part of the semi-solid slurry. When the semi-solid material is injected under pressure into the mould, the liquid part solidifies around the grains of the preexisting α phase forming a fine microstructure of secondary α-Mg grains and eutectic phases. The microstructures of the ZEC alloy in the as-cast state are shown in figures 1 and 2, while the microstructures of the MEZ alloy are shown in figures 3 and 4. In the case of the ZEC alloy, two phases were detected around the α-Mg grains. Intergranular precipitates containing Mg-Ca-Zn as well as eutectics containing Mg-Zn-Ca-RE were found decorating the border of the primary α-Mg grains. The precipitates are possibly of the type Mg2X, where X = Ca, Zn and the eutectics are possibly of the type Mg6Zn3Ca2, where atoms of Zn and Ca are substituted by atoms of Ce, Nd and La. The phases on the ZEC alloy are named after the ones founded in the studied Mg-Ca-Zn alloys [2, 4] because of the similarities on the atomic compositions reported by EDX.

70

Figure 1. Microstructure of the thixocast ZEC alloy, BSE-SEMEM

Figure 2. Detail of the microstructure of the thixocast ZEC alloy, BSE-S

Figure 3. Microstructure of the thixocast MEZ alloy, BSE-SEM

Figure 4. Detail of the microstructure of the thixocast MEZ alloy, BSE-SEM

In the MEZ alloy, grain boundary intermetallics containing Mg-RE-Zn were clearly identified; these intermetallics correspond to the well known type Mg12X, where X = RE, Zn, [3, 5, 6]. After the comparison of the figures 1 and 3 it is clear that the microstructure of the ZEC alloy is finer than the microstructure of the MEZ alloy. This fact may be explained because of the Ca additions on the ZEC alloy, which is known as a microstructure refiner in magnesium alloys [4]. EDX analysis on the α-Mg grains of the ZEC alloy revealed 0.3 % maximum of Ca content. On the MEZ alloy, 0.3 % maximum of Ce was detected. No coring effect was found on both alloys, as expected in a semi-solid process. 3.2

Mechanical behaviour

The ultimate tensile strength and the total percent elongation of the thixocast alloys ZEC and MEZ are displayed in table 2. These values are reported together with those found in the literature for the high pressure die casting MEZ [2]. The values of UTS are slightly higher for the ZEC alloy than that of the MEZ alloy; this fact is not surprising and it is attributed to the beneficial effect of Ca additions on improving the mechanical strength of Mg alloys [2, 4]. Although the UTS value of the HPDC

71 MEZ alloy is 20% higher than that obtained for the semi-solid MEZ, the elongation values are in good accordance for both processes. Table 2. Tensile properties of the semi-solid ZEC and MEZ alloys

0.2 % YS (MPa) 89.4 73.8 97.6

ZEC SS MEZ SS MEZ HPDC

UTS (MPa) 116.5 111.2 134.8

El. (%) 1.7 2.8 2.9

12

200°C

Creep strain, %

10

8

ZEC σ = 80 MPa

6

4

2

175°C 150°C

0

0

25

50

75

100

135°C

125

150

Time, h

Figure 5. Creep curves for thixocast ZEC at 80MPa 12

200°C

Creep strain, %

10

8

MEZ σ = 80 MPa

6

4

2

0

0

25

50

75

175°C

135°C

150°C

100

125

150

Time, h

Figure 6. Creep curves for thixocast MEZ at 80Mpa

Typical creep strain versus time curves of the semi-solid alloys at 80 MPa are presented in figures 5 and 6. The ZEC alloy exhibits always a lower primary creep strain than that of the MEZ alloy, which is related to the better mechanical strength achieved by effect of Ca additions, but a higher minimum creep rate, particularly at 175 and 200 °C. A 12 % creep strain was reached for the ZEC alloy in 71 hours at 200 °C, whereas for the MEZ alloy 145 hours were needed.

ZEC 60 MPa ZEC 70 MPa ZEC 80 MPa MEZ 60 MPa MEZ 70 MPa MEZ 80 MPa

10

8

70h

Creep extension at 100 h, %

12

70h

72

6

4

2

0

135°C

150°C

175°C

200°C

Figure 7. Total creep extension of the thixocast alloys after 100 hours of testing

As can be seen in figure 7, the total creep strain accumulation after 100 hours is lower for the ZEC alloy at 135, 150 and 175 °C, that is, primary creep controls the total creep extension. At 200°C the total creep extension of the ZEC alloy surpasses that of the MEZ alloy. Because of physical limitations of the extensometers on measuring creep strains higher than 12 %, the tests at 200 °C for the ZEC alloy at 70 and 80 MPa were stopped after 70 hours. After 200 hours of testing, the total creep extension of the ZEC alloy at 135 and 150 °C is still lower than that of the MEZ alloy, controlled by the primary creep strain. However, at 175 °C and 80 MPa, the total creep strain of the ZEC alloy surpasses that of the MEZ alloy after 125 hours, as can be observed in figure 8. 2.5

Creep strain, %

2.0

MEZ

1.5

ZEC T = 175°C σ = 80 MPa

1.0

0.5

0.0

0

25

50

75

100

125

150

Time, h

Figure 8. Creep curves for the studied thixocast alloys at 175°C and 80 MPa

Minimal secondary creep rates of the semi-solid alloys at different temperatures and stresses are presented qualitatively in figure 9. At all temperatures and stresses, the secondary creep rate of the ZEC alloy is higher than that for the MEZ alloy. However, the differences on creep rates are higher at 175 and 200 °C.

73 1E-7

Log εs

1E-8

ZEC 60 MPa ZEC 70 MPa ZEC 80 MPa MEZ 60 MPa MEZ 70 MPa MEZ 80 MPa

1E-9

135°C

150°C

175°C

200°C

Figure 9. Minimal creep rates (s ) of the semi-solid alloys at different temperatures and stresses -1

4

Conclusions

The thixocasting forming process was successfully applied on Mg-Zn alloys without Al content. Net-shaped pieces with minimal porosity and globular microstructure were obtained. Calcium has a beneficial effect on the mechanical strength of Mg-Zn alloys. Nevertheless, the effect of Ca on the creep properties of Mg-Zn-RE alloys is still under study. The proposed ZEC alloy presented good creep properties at 135 and 150 °C in comparison with the high temperature creep resistant MEZ alloy, with high content of RE. The values of n and Q are on progress and are not presented in this study because additional creep tests are required to understand the creep mechanisms at low (50-70 MPa) and high stresses (80-100 MPa).

5

References

S. Kleiner, O. Beffort, A. Wahlen, P.J. Uggowitzer, J. Light Metals 2002, 2, 277-280 T.V. Larionova, W.-W. Park, B.-S. You, Scripta Mater. 2001, 45, 7-12. D. Weiss, A.A. Kaya, E. Aghion, D. Eliezer, J. Mat. Sci. 2002, 37, 5371-5376 A. Finkel, L. Shepeleva, M. Bamberger, E. Rabkin in Magnesium Technology 2002 (Ed.: H.I. Kaplan), TMS, USA, 2002. [5] I.P. Moreno, T.K. Nandy, J.W. Jones, J.E. Allison, T.M. Pollock, Scripta Mater. 2003, 48, 1029-1034 [6] F. Villa Fanjul, S. Srimanosaowapak, K.R. McNee, G.W. Greenwood, H. Jones, Z. Metallkunde 2003, 94, 25-29 [7] T. Miyazaki, J. Kaneko, M. Sugamata, Mater. Sci. Eng. 1994, A181/A18, 14101414.

[1] [2] [3] [4]

74

Creep Resistant Magnesium Alloys for Powertrain Applications Mihriban. O. Pekguleryuz McGill University, Department of Metals and Materials Engineering 3601 University Street, Montreal, QC, Canada H3A 2B2

A. Arslan Kaya Tubitak Marmara Research Center, P.O.B.21, Gebze, 41470 Turkey

Abstract Creep resistant magnesium alloys are candidate materials for automotive powertrain applications. Since the 90’s, a number of new creep-resistant magnesium alloy systems have been investigated and developed. These are for the most part based on rare-earth, alkaline earth, and silicon additions. This paper gives an overview of creep resistance in magnesium and a review of creep resistant magnesium alloys for power-train applications.

1

Introduction

Automotive use of magnesium is currently limited to a few select applications such as the instrument panel, steering wheel, and valve cover. The alloys used in these applications utilize Mg-Al based casting alloys such as the AM series, Mg-(2-6)%Al and the AZ91 alloy (Mg-9%Al-1%Zn). These alloys offer good combination of room-temperature strength and ductility, good salt-spray corrosion resistance and excellent diecastability. Major growth areas in automotive use for Mg alloys are the powertrain applications, namely, the transmission case and the engine block. These applications require good elevated temperature performance at service conditions of 150-200 °C and stresses in the range of 50-70 MPa. Commercial AM and AZ type of magnesium alloys do not possess this required elevated temperature performance. In the last decade, concentrated effort has been expanded in developing and optimizing magnesium alloys with good creep resistance. Mg alloys for high temperature applications require considerations on several properties in conjunction with the creep strength. Castability, especially die-castability, of a new alloy must be at an acceptable level, as this is the major processing route for production of parts made of Mg alloys. Corrosion and oxidation resistance is important. In addition, to be able to compete with relatively lower cost metals such as aluminum and steel, the alloys should not have additional cost penalty by containing costly alloying elements. Majority of magnesium parts used in automotive applications is high-pressure diecast. This is due to the excellent fluidity of magnesium alloys providing good die filling, the compatibility of liquid magnesium with die materials resulting in long die-life and the high pro-

75 ductivity of the die casting process upsetting the material cost factors. However, gravity casting is also beginning to attract interest especially for higher performance applications such as the engine block. In this review, background on creep resistance of magnesium alloys is presented. Creep resistant magnesium alloys are classified into two groups; (1) Creep resistant magnesium alloys used for high-pressure die-casting and (2) other creep resistant magnesium alloys used for low-pressure die-casting and gravity casting.

2

Creep Deformation

Creep is time-dependent slow and continuous deformation in time. While in elasticplastic deformation strain is a function of stress only, in creep strain is a function of stress, time and temperature.

ε = f(σ) elastic/plastic deformation, ε = f(σ, T, t) creep deformation Creep deformation is due to the thermally activated (short-range) component of the stress,

σ = σ E + σ* (normal stress) τ = τ u + τ* (shear stress) where σ*, τ* are thermal and σE, τ u are athermal components. The athermal component involves movement of dislocations passed stress fields whose origins are distant on an atomic scale, about 300 atomic distances. This long-range component is associated with work hardening. The thermal component is due to atom vibrations. It involves dislocation motion passed obstacles whose stress fields are more localized. Hence, while the long-range component involves the movement of long segments of parallel dislocations, the thermal component can only cause movement of small sections of a dislocation (e.g. dislocation jogs, or the cutting of a dislocation by another). This nevertheless causes slow and permanent deformation called creep. Since atoms vibrate anywhere above absolute zero, thermal component exists at very low temperatures and creep is possible even at very low temperatures, given sufficient time. Fig. 1 shows a typical creep-strain versus time curve for a metal, where three distinct regions can be observed. After an initial and instantaneous restructuring of the metal, primary stage creep occurs with increasing creep rate. Here the work hardening is higher than the stress recovery. In the secondary creep (steady-state) stage the metal experiences a threeway balance between work-hardening, increasing stress and stress recovery Fig. 1. Typical creep strain vs. time curve showing the resulting in steady-state creep (conthree stages of creep stant creep-rate). The tertiary creep stage exhibits increasing creep rate due to necking resulting finally in fracture.

76 -1

10-1 10

C onventional plastic flow

έ = Aσ

n

έ /R T ) e (-Q

-2

σ/G

10-2 10

Y ield strength -3 -3

10 10

-4

10-4 10

C ore diffusion D islo catio n creep B ulk diffusion D iffusio nal creep

έ = B σ e (-Q /R T )

E lastic deform ation only

Boundary diffusion -5 -5

10 10

0

0.5

T /T M

έ

σ

Bulk diffusion

1.

Mathematical description of creep that considers a thermally activated process capable of acting in the direction of the stress or opposite to it as it passes over an energy barrier defines creep-rate as shown below (1). Hence, creep rate at a given stress is extremely temperature sensitive. έ = A e (-q/kT) ⋅ e(ντ*/kT)

Normalized stress vs. homologous temperature maps (Fig. 2) show the Fig. 2. Normalized stress vs homologous temperature two distinct creep regions, (1) dislocamaps showing the various regions of deformation (2) tion creep at high stresses where strain-rate has a power law dependence on stress and (2) diffusional creep at high temperatures with a linear dependence of strain-rate on stress. The mechanisms of creep at micro-level are recovery mechanisms as summarized in Table I. Table 1. Creep mechanisms (1)

Low to Moderate Temperatures: High temperatures : Thermally Activated Recovery Stress-activated Recovery Dislocation Creep Diffusion creep • Grain sub structures. In HCP structures like Mg these are elongated and parallel to grain boundaries, in FCC they are equiaxed (Fig. 3.a). • In HCP Mg, twinning (Fig. 3b).

(a)

(b)

• Dislocation Intersection • Fast diffusion paths such • Dislocation climb as dislocation cores and • Movement of dislocation grain boundaries atmospheres • Bulk diffusion (vacancy • Cross-slip and interstitial • Grain-boundary shear • T > 0.3TM

(c)

Fig. 3. (a) sub-grain formation in Mg (b) twinning and basal slip in Mg. (c) non-basal slip concentrated at grain boundaries rather than grain interiors at high temperatures (Ref. 3).

Magnesium, having h.c.p.-crystal structure, possesses fewer slip systems compared to, for example, f.c.c structure, and therefore its ability to deform is limited especially at low temperatures. The slip system at the room temperature is basal{0001} 〈1 120〉 and above 250°C, {1011} and {10 12} pyramidal and {10 10} prism slips are activated. Room tempera-

77 ture deformation occurs mainly by slip on the basal planes. In addition, twinning on the pyramidal {10 12} planes provides means of deformation Mechanisms of creep in pure Mg have been studied (4) and it has been observed that pure Mg exhibits basal slip on the {0001} planes, twinning and sub-grain formation (Fig. 3.a&b) in the 90-300 °C temperature range at stress levels of 8-70 MPa. These are mechanisms in the primary creep stage. At the higher temperatures, above 250 °C, non-basal slip on the pyramidal and prismatic slip planes (Fig. 3c) and grain boundary deformation and sliding occur (steady-state creep). Creep mechanisms in a metal or alloy can be identified by determining the activation energy for creep from Arrhenius relationship of strain rate and reciprocal temperature. Minimum creep rate in steady-state (secondary) creep is obtained from long-term creep tests. The value of the activation energy and the stress exponent would give clues as to which thermally activated creep mechanisms listed in Table II, are predominant in a particular temperature and stress range. Table 2. Strain-rate and activation energy for thermally activated creep mechanisms (1)

Creep Mechanism

Creep Rate / Activation Energy

Dislocation Intersection - work required to force a disloca- έ = A e [-(qi - τ*b2l)/kT] tion through the stress field of another can be thermally aided. Mg single crystals Activated Cross Slip - Thermal activation to unite partial q =100 kJ/mol dislocations and to breakdown into partials on different slip for high stacking fault systems. metals. 2 Movements of Dislocations with Jogs - formation (qf) and έ = A e [-(qd - τ*b x)/kT] movement (qm) qd Mg 135 kJ/mol, of vacancies. Self diffusion (qd = qf + qm). qMg 125 kJ/mol. α έ = Aσ e (- qd ) / kT Dislocation Climb - Climb of dislocations over sessile βσ dislocations or precipitates. Vacancy diffusion to and from έ = Be e (- qd ) / kT edge dislocations (qd) and formation of jogs (qj). Activation q in the range of qd (135 kJ/mol) energy, q = qd + qj Movement of Dislocation Atmospheres - Diffusion of solute έ = A σ e (- qs ) / kT qAl in Mg 143 kJ/mole atoms and the viscous behavior of the solute atmosphere Grain Boundary (GB) Shear - relative movement of grains, deformation in very narrow region adjacent to the GB, the shear direction lies in the boundary with the max resolved shear stress. Affective at high temperatures, usually when the recovery temperature is achieved. Discontinuous. GBs recover before the grains, softening. Accounts for 30% of the deformation. Grain boundary diffusion: 80-100 kJ/mol. Creep mechanism in AZ91D (Mg-9Al-1Zn) alloy has been studied by Dargusch et al (5). It was found that in the 125-175°C temperature range and with σ = 50 MPa, the n activation energy, Q in έ = Aσ exp (-Q/RT) was 30-45 kJ/mol and n = 2. This activation energy was too small compared to self-diffusion of Mg (135 kJ/mol), diffusion of Al in Mg (143 kJ/mol) or grain boundary diffusion (80 kJ/mol). The value did fit with the activation energy Q for discontinuous precipitation of Mg17Al12 = 30 kJ/mol. Creep induced precipitation of Mg17Al12 and grain boundary migration were observed in the evolved

78 microstructure after creep. Since such precipitation would aid easy grain boundary sliding and migration, the creep was attributed to creep-induced precipitation of Mg17Al12 and the involvement of this phase in the creep deformation of Mg-Al based alloys was established.

(a)

(b)

Fig. 4. (a) Creep induced precipitates of Mg17Al12 and (b) grain boundary migration in AZ91D (Ref. 5).

At stresses higher than 50 MPa, Q was determined to be 95 kJ/mol and n was 5, indicating a change in the creep mechanism. No mechanism was named in the study for the high stress range but looking at Table II this activation energy seems to fall in the range of activated cross slip, as well as grain boundary diffusion for magnesium. Other potential creep mechanisms would be dislocation intersection or overcoming of Peierls stresses. Creep mechanism in AM50 (Mg-5% Al) alloy has been studied by Terada et al (6). At 100 MPa, and in the temperature range of 150-225 °C, the activation energy was determined to be 126 kJ/mol, which the authors attributed to dislocation climb. It would fit very well with a creep mechanism described as “movement of dislocations with jogs” (Table II).

3

Creep Resistant Magnesium Alloys for High Pressure Diecasting

3.1

Mg-Al-Si Alloys

The first development in creep resistant magnesium diecasting alloys for automotive applications has been the development of Mg-Al-Si diecasting alloys by Volkswagen in the 70’s. Silicon addition in AS alloys leads to the formation of the Mg2Si precipitates that has low density, high hardness, low thermal expansion coefficient, and perhaps most importantly, unlike the β-Mg17Al12 phase, a high melting temperature (1085 °C). Si lacks the ability to tie up Al rendering it ineffective in hindering the formation of β-phase and discontinuous precipitation for a given level of aluminum. These alloys hence have two types of intermetallics, the β- (Mg17Al12), and the Mg2Si phase with high temperature stability, which is deemed to pin dislocations and the grain boundaries thereby imparting increased creep resistance to the alloy. The Mg2Si that has a face centered cubic crystal structure (7) and is a metallurgically stable phase that forms in Chinese script morphology under slow cooling conditions, which would exhibit very low ductility. In the diecast alloy, a fine distribution of the strengthening phase, Mg2Si, is achieved with resultant high temperature resistance and a satisfactory level of toughness (8). The creep resistance achieved is moderate and the alloys are difficult to diecast.

79 The AS41 (Mg-4Al-1Si) and AS21 (Mg-2Al-1Si) alloys are the two standardized Mg-Al based heat-resistant alloys in use that display some improvement in creep properties in the range 130-150 °C. AS41 offers better heat and wear resistance than AZ91, albeit with reduced fluidity of its melt. AS41 has moderate castability, but due to lower Al content, castability in AS21 becomes even a bigger problem. Foerster (9) suggested that silicon improves the fluidity of a magnesium alloy when 4%Al is present, making AS41 more castable than AS21. Considerable amount of data on Mg2Si has been reported by Beer et al. (7). However, silicon addition must be accompanied by a reduction in aluminum content to lower the amount of Mg17Al12 in the microstructure in order to improve the creep resistance. Creep mechanism in AS21 alloy was studied by by Dargusch and Dunlop (5, 10). The same activation energy of 30-45 kJ/mol and n = 2 were determined and creep induced precipitation of Mg17Al12 was observed (Fig. 5). The improved creep resistance of the AS21 over AZ91D was attributed to the smaller amount of the creepinduced precipitate and the grain boundary pinning provided by the Mg2Si precipitates. Blum et al. (11) have reported that at high stresses AS21 displays smaller maximum deformation resistance compared to AZ91 whereas Fig. 5. Microstructure of AS21 after 120 hr under low stress the opposite is true. It is also creep deformation at 150 C and 50 MPa interesting to note, as reported by Aune and Ruden showing the Mg2Si and creep induced (12), that up to 200 °C, alloys of the AZ, and AS, precipitates of Mg17Al12 (Ref. 5). series all appear to exhibit the same relative decrease in strength as the temperature is increased, and that at all temperatures AZ91 shows superior yield strength. The same authors also revealed that as 200 °C is approached except for AS41 these alloys would display similar levels of ductility. It is possible to obtain a compromise between castability and creep resistance in this alloy system by going to an AS31 (Mg-3Al-1Si) composition. Recently, a modified version of AS21x with small additions of rare earth elements has been developed. Rare earth additions at trace levels do not affect the creep resistance but improve the salt-spray corrosion resistance as well as castability (13). Mg17Al12

3.2

Mg-Al-RE Alloys

A major development in creep resistant magnesium alloys has been the emergence of rareearth (RE) containing alloys (AE41, AE42). This group of Mg-Al-RE alloys involves at least one, and in general a mixture, of the REs as precipitate forming alloying addition in their compositions. Rare earth alloying additions are expensive and therefore a cheaper substitute known as mischmetal, a mixture of several rare earth elements that is enriched in one of the constituents is generally used. AE series constitutes the standardized alloys in the group of Mg-Al-RE alloys. This alloy system exhibits major improvement in creep resistance due to the complete suppression of the formation of the β-(Mg17Al12) phase and the presence, instead, of the Al-RE containing intermetallics (15). In a review by Aune and Westengen on mechanical properties of various magnesium alloys, it has been shown that the ultimate tensile strength increases with the Al and RE content in the alloy, and that tensile yield strength primarily depends on the amount of RE and attains a limiting

80 maximum level (14). For a given rare earth content in the alloy the creep strength was reported to decrease with increasing aluminum content, while for a given aluminum content it increased with increasing amount of rare earth addition (15). There have been a number of studies aimed at identifying the second phases in AE alloy family. Polmear has claimed that Mg-Al alloys containing RE addition are suitable only for diecasting since slow cooling rates lead to the formation of coarse Al2RE particles (16). Mg12Ce type particles were also detected at the grain boundaries during creep. Both of these precipitates, unless they coarsen, are considered to inhibit grain boundary sliding and thus being beneficial for creep resistance (16). Precipitation of Al4RE compound in the alloy AE42 has also been reported (5, 17-19). Depending on the composition, precipitates of other types of compounds have also been mentioned in the literature. Pettersen et al (17) suggested that when RE/Al weight ratio is above 1.4 all of the aluminum will be tied up as Al11RE3 in which case further precipitation of other phases such as another type of Al-Re phase or Mg12RE becomes possible. In samples produced by cold chamber diecasting they observed Al11RE3 and Al10RE2Mn7 formation at the grain boundaries. Their TEM observation has revealed that the major phase was Al11RE3 (a body-centered orthorombic cell with lattice parameters a = 4.5 Å b = 13.2 Å and c = 9.9 Å). No special orientation relationship was found between this phase and Mg-matrix. EDS measurements gave an atomic ratio Al:RE = 3.3 ± 0.6 which is claimed to be in good agreement with the proposed stoichiometric ratio 11 :3. The relative amounts of REs were reported to be (at%) La27±2Ce51±2Pr6±1Nd16±1. This phase was found to co-exist with a minor phase Al10RE2Mn7. The diffraction patterns from this phase were indexed according to a hexagonal system with lattice parameters a = 9.0 Å and c = 13.1 Å. The extinctions observed show that it is a trigonal phase as only reflections with –h + k + l = 3n are allowed. This structure was found in agreement with the published data for Al10Ce2Mn7 phase. There was, however, a problem due to an unexpected reflection, (1011); several pictures indicated that –h + k – l = 3n reflections were also allowed which is not consistent with any space group. The authors explained this as being due to a mirror image of the reciprocal space with the basal plane as the mirror. A possible physical interpretation given is that growth of twins with the basal planes as the twinning-plane may be influenced by the substitution of RE for Ce. EDS analysis gave Al49±2RE12±2Mn39±2 which is in reasonable agreement with the proposed 10:2:7 ratio. The relative compositions of the RE components was La10±1Ce50±3Pr9±1Nd31±2. Lanthanum was said to go into Al11RE3 phase, while Nd into Al10RE2Mn7 phase. They attributed the creep resistance to the precipitate morphology. Since the precipitates forming at the grain boundaries have a strongly branched morphology anchoring into several neighboring grains together, this was deemed to contribute to grain boundary pinning. Dargusch et al have studied creep in AE42 alloy (5) and have observed that just as AZ91 and AS21 alloy, the activation energy was in the range of 30-40 kJ/mol, indicating the same creep mechanism of creep-induced-precipitation resulting in grain-boundary migration. Mg17Al12 was not observed but Al-supersaturation was noted. The difference between AZ and As alloys and the AE was that the onset of the second creep mechanism at the high stress region at 150 °C was not seen in AE42 up to 100 MPa and that the creep strain was lower than AZ91 and AS21 (Fig. 6). In a more recent paper by Powell et al (20) the microstructural evolution of AE42 alloy was studied (Fig. 7). It was reported that two different intermetallic phases Al11RE, Al2RE and practically no Mg17Al12 were present in the diecast structure. Creep testing

81 at 150 °C and 175 °C showed a change in microstructure and changes in the relative amounts of these phases (Table III). These changes are consistent with the reaction by which Al11RE3 decomposes to Al2RE and frees Al, which then reacts with Mg to form Mg17Al12, which is accepted as being responsible for poor creep resistance in magnesium alloys (21). To see if stress played a role on the decomposition of Al11RE3 Fig. 6. (a) Creep strain vs time and (b) creep rate vs stress samples were exposed to the same in AE42, AS21 and AZ91 (Ref 5). thermal history but without the applied stress (20). The authors observed the same microstructural changes and concluded that Al11RE3 suffered merely from thermal instability at 175 °C. A relationship was also discovered between the relative amounts of Al11RE3 vs. Al2RE and the La:Nd. (below about 0.7 Al2RE was seen and above 0.7 Al11RE3 seemed to form). They indicated that the formation of either Al11RE3 or Al2RE is sensitive to individual rare earth elements, and that this might influence the stabilization of the Al11RE3 phase in AE42 alloy. 150 C, 50 MPa

150 C,

Table 3. Change in Amounts of phases (wt %) with creep test in AE42 alloy (20).

T (C) Mg Al11RE3 Al2RE Mg17Al12 As 97.5 1.8 cast 25C 150 97.7 1.5 175 97.0 1.2

(a)

0.8

0.0

0.8 1.3

0.0 0.6

(b)

Fig. 7. Microstructure of AE42 (a) diecast (b) after 175 °C exposure (20).

AE42 alloy was perceived to have issues with cost, castability and fatigue strength and was not considered for wide-spread automotive use for a long time. In recent studies a fresh look at the AE alloy family is presented (22). Different combinations of Al and RE are investigated namely Mg-3%-5%RE (AE35), and Mg-8Al-2RE (AE82). Future studies on AE42 should address the thermal stability at 175 °C and conduct a thorough search in the ternary Mg-Al-RE alloy system to determine compositions that offer both creep resistance, thermal stability and good castability. Phase diagram studies in the ternary system are also important. 3.3

Mg-Al-Ca Alloys

Ca addition, as a cheaper and lighter alternative to RE elements, also contributes to high temperature properties (23-28). When Ca is added to Mg-Al binary alloys, the type of precipitating compound is found to depend on the Ca/Al mass ratio. When this ratio is more than ~0.8 presence of both Mg2Ca and Al2Ca were detected that resulted in considerable increase in hardness, while below this ratio only Al2Ca was observed to have formed (Figs. 8 & 9). Both types of precipitates were observed to form along the grain boundaries with an increasing volume fraction as the Al and Ca content increased (27).

82

Fig.8. Diecast microstructure of Mg-5Al-0.8Ca (Ref. 28).

3000

3000 2400

Pekguleryuz et al have shown that the creep resistance in Mg-Al-Ca ternary alloys (AX alloys) in the range of 2-6 wt% Al and 0.61.0wt% Ca show creep resistance at 150 °C similar to the AE42 alloy with the added benefit of good corrosion resistance (25, 28, 29). The Al2Ca compound has stability at high temperatures and the formation of the βMg17Al12 phase is suppressed. Beneficial effect of Ca addition was also proven in rap idly solidified AZ91 alloy. Gjestland et al (29) have shown that despite the smaller grain size com-

Al 2Ca

Al2Ca

2400

Mg2Ca 1800

1200 / SEC

1200 CNTS / SEC

1800

CNTS 600 0

(a)

15 55

25

65

35

45

275 THETA

600

0

15

25

35

45

55

65

75

2 THETA

(b)

Fig. 9. XRD from (a) Mg-5Al-0.6Ca alloy showing Al2Ca. (b) Mg-5Al-0.8Ca showing the presence of Al2Ca and Mg2Ca.

pared to conventional AZ91, rapidly solidified structure involving fine Al2Ca particles exhibited drastically smaller secondary creep rate. Although useful, Ca addition is not problem-free. Calcium makes the alloy prone to hot cracking and die-soldering especially when the level is above 1% (23-25, 30). Subsequent addition of 8% Zn to the alloy has improved diecastability but the properties were found to vary over a wide range (31). The second phase in this alloy was a Mg-Al-Ca-Zn containing complex phase. At this time, none of these experimental alloys with Ca addition has yet led to a commercially successful alloy, More work with the base alloy systems containing Ca is needed to optimize these alloys for diecasting, however, work has been carried out to use these alloys for semi-solid applications (32). Mg-Al-Ca system can be further investigated for optimum castability. Recently, Luo (33) has reported that increasing the Ca to above 2% improves castability by eliminating hot-tearing. Terada (6) et al studied creep mechanisms in Mg-5Al-1.7Ca alloy. At 150 °C, the change in the stress exponent occurred at around 120 MPa, and in the low stress regime (100 MPa), in the 175-250 °C range, the activation energy was determined to be 142 kJ/mol. While the authors attributed this to dislocation climb, the mechanism can be better described by the mechanism of “Movement of Dislocation Atmospheres” (diffusion of solute atoms and the viscous behavior of the solute atmosphere) since q is identical to qAl 143 kJ/mole (diffusion of Al in Mg) (Table II).

83 3.4

Mg-Al-Sr System

Mg-Al-Sr (AJ) alloys are a new addition to the creep-resistant magnesium alloys (34-43). Various alloy compositions such as AJ51 (Mg-5Al-1Sr), AJ52 (Mg-5Al-2Sr) and AJ62 (Mg-6Al-2Sr) have been studied. AJ52 shows the highest creep resistance, while AJ62x gives an excellent combination of creep performance and castability. Pekguleryuz et al report that the alloys show different microstructures based on the Sr/Al ratio (42). For Sr/Al ratio below about 0.3, Al4Sr intermetallic is the only second phase in the structure (Fig. 10a&b). When the Sr/Al ratio is higher, a second intermetallic phase (a ternary Mg-Al-Sr phase) is also observed. Sr/Al controls the formation of Mg17Al12 as well. When the Sr/Al ratio is very low, there would be insufficient amount of Sr to bind all Al and the excess Al would form the Mg17Al12 phase. It is reported that the ternary Mg-Al-Sr phase is a new compound, which is under study. The superior creep resistance of AJ52 and AJ62 alloys is attributed to the existence of this second phase in addition to Al4Sr. Al4Sr is a tetragonal phase and no phase coherency with the primary magnesium phase has yet been reported. Future work in this alloy system should focus on the determination of the ternary phase diagram and the identification of the ternary phases.

Al 4Sr Al4Sr

Mg 17 Al 12

(a)

(b) Al4Sr

Mg-Al-Sr phase

Mg-Al-Sr phase

Al4Sr

(c)

(d)

Fig.10. Diecast microstructures of (a) AJ51(Mg-5Al-1Sr) and (b) AJ62L (Mg-6Sr-1.6Sr) alloys with the Sr/Al ratio is lower than 0.3, exhibiting AL4Sr second phase. Microstructures of (c) AJ52 (Mg-5Al-2Sr) and (d) AJ62 (Mg-6Al-2Sr) alloys with the Sr/Al ratio higher than 0.3 exhibiting two types of second phases, Al4Sr and a ternary Mg-Al-Sr compound.

84 3.5

Magnesium Diecasting Alloys with Combined Additions of Alkaline Earth and/or Rare Earth Elements In recent years, many Al-containing Mg alloy systems have been developed with combined additions of rare-earth and alkaline earth elements. Nissan patent (44) on a Mg-Al-Ca-RE alloy and later a Honda alloy (45) ACM522 (Mg-5Al-2Ca2RE) both claim improved creep resistance over AE42 alloy. ACM522 alloy is based on specific AI/RE/Ca ratios to ensure AI-RE precipitates in addition to Al-Ca intermetallics. This alloy is used by Honda in oil pans for their hybrid car.

Fig. 11. Microstructure of Mg-8Al-2.5RE1.6Ca-1.3Mn (permanent-mold cast)

The microstructures of these alloys exhibit a combination of Al-RE and Al-Ca intermetallic mixtures. The permanent-mold cast microstructure of Mg-8Al-2.5RE-1.6Ca1.3Mn version of the alloy given in Fig. 11, shows Al2Ca coupled eutectic phase as well as needle shaped Al-RE (Ca, Mn) intermetallics and Al-Mn-RE intermetallics (46). Another alloy system based on Mg-Al-Ca-RE with optional additions of Sr and Zn is a patent by Dead Sea Magnesium and Volkswagen AG (47). Two alloy formulations codenamed MRI 153 and MRI 230 are reported to exhibit good elevated-temperature performance (48) at 150°C and 180°C respectively. Furthermore MRI 153 exhibits good diecastability similar to AZ91 alloy. The alloy system of this patent is quite complex, having a number of elements and would lend itself to further development to find optimum combinations of strength, creep resistance and castability. Two separate alloy systems with combined additions of Sr and Ca but no rare-earths have been developed by Noranda (49) and General Motors (50). The Noranda alloy is a Mg-Al-SrCa alloy with small amounts of Sr and Ca (AJX Alloy). Code-named as Noranda N alloy, it has creep resistance comparable to AS41 at 150°C but creep resistance improves at 175°C being superior to AE42 alloy (46). The GM version is a Mg-Al-Ca-Sr alloy with substantial Ca and a small amount of Sr (AXJ Alloy) and exhibits superior creep resistance (51). All these complex alloys offer great potential for high creep resistance and the major effort should be invested on improving diecastability of these alloy systems compositions that offer the highest creep resistance. This can be achieved by systematic and fundamental castability research or by trace element effects. 3.6

Comparative Properties of Magnesium Diecasting Alloys for Powertrain Applications

Formal comparative evaluation of the various alloys is in its initial stages. Automotive partnership programs such as USCAR and EUCAR have projects that evaluate the various creep resistant magnesium alloys. Important performance requirements for powertrain components, among others, are tensile and compressive creep resistance, bolt-loadretention, yield strength, fatigue resistance, corrosion resistance and diecastability as well as cost and availability. It is rather difficult to make a rigorous comparison of the alloys from published data due to differences in sample, test conditions and type of test. Creep tests are carried out at

85 different stress-temperature conditions and results are reported at different creep time. An attempt is made in Tables IV & V to summarize some of the data on the various alloys. Other data can be found in a recent review by A. Luo (52). Formal evaluation programs would be of great benefit to end users in identifying appropriate alloys for powertrain application development programs. Table 4. Tensile properties (42, 48, 52)

ALLOY 20°C

150°C

UTS, YS, MPa MPa AZ91D 204 (240) 139 (160) AS41 197 (215) 136 (140) AS21x* 210 121 AE42 220 (230) 128 (145) AX51 192 128 AX52 228 161 ZAX850 219 146 AJ52x 212 134 AJ62x 240 143 AJ62Lx 276 153 AJX500 236 133 AXJ531 238 190 ACM522 200 158 MRI15* 197 (250) 157 (165) MRI230D 235 180 A380 290 (290) 155 (180)

%E 3.1 (3) 4 (6) 5.5 10 (11) 7 13 5 6 7 12 10 8 4 2.2 (5) 5 3.2 (3)

UTS, MPa 169 (157) 153 (155) 130 140 (157) 161 159 163 166 163 152 175 205 251 (257)

175°C YS, MPa %E

UTS, MPa 105 (108) 16 (23) 138 94 (95) 17 (24) 127 (127) 87 20 110 88 (100) 23 (20) 121 (130) 102 7 171 117 11 110 12 141 116 27 143 108 19 143 101 16 137 196 138 7 152 118 139 150 16 148 (159) 6.4 (7) 248 (228)

YS, MPa %E 89 85 (85) 78 81 (91) 133

21 19 (25) 23 23 (23) 23

100 18 103 19 109 27 98 20 146 15 132 9 113 (110) 3.4 154 (148)7.1 (10)

Table 5. Creep of diecast alloys (42, 48, 52)

ALLOY TENSILE CREEP, % @ 50 MPa, 200 hrs AZ91D AS41 AS21x AE42 AX52 AXJ531 AJ52x AJ62x AJ62Lx MRI 153 MRI 230D A380

@ 50 MPa, 500 hrs

70 MPa, 100 hrs

70 Mpa, 200 hrs (53)

% Compressive Creep 150 °C, 70 MPa, 200 hrs

150 °C 2.7 0.05 0.19 0.06 0.04 0.05 0.13 0.18 -

175 °C * 2.48 1.27 0.33 0.05 0.05 0.29 -

150 °C 6.35 0.07 0.08 0.03 -

175 °C 0.44 0.09 -

175 °C 0.06 0.06 -

175 °C 8.95 0.18 0.14 0.24

21 6.13 3.97 2.16 0.24 1.73 -

0.08

0.04

0.10

0.05

-

0.22**

0.03

*failed after 80 hrs

**A383 alloy

86 Diecastability is a special performance requirement of these alloys. The factors that contribute to good diecastabilty are low surface tension and relatively low liquidus for adequate fluidity, resistance to hot-tearing, optimum degree of wettability to avoid excessive reaction with steel dies, etc. Generally increasing aluminum in Mg-Al based alloys improves diecastability by decreasing the liquidus and improving fluidity. The same can be achieved by increasing the diecasting temperature, within reasonable limits so as not to sacrifice die-life. A larger freezing range may also help diecastability by increasing the length of time the metal stays liquid and hence eliminate cold-shuts. Alloying elements that decrease the thermal conductivity or increase heat capacity are also beneficial for diecasting (40). Each alloy would require special set of diecasting parameters to fit the specific characteristics of the alloy.

4

Creep Resistant Magnesium Alloys for Gravity or Low Pressure Casting

Traditionally automotive use of magnesium has been primarily in die cast parts to take advantage of the high productivity of the die casting process that balanced the relatively high cost of magnesium alloys. With the price trends of magnesium since 2001, magnesium gravity cast or low-pressure parts are also becoming cost-effective. Development of gravity cast alloys does not pose the additional challenge of meeting diecastability requirements. This would theoretically allow for a wider selection of alloy systems and compositions to be explored. Despite the low creep resistance of pure magnesium, it is known that alloying has been quite successful in the past in developing magnesium alloys for high temperature service for aerospace applications. This is evident in comparing the “maximum usefulness temperature” of various metals. A measure of maximum usefulness temperature is arbitrarily defined as the temperature at which the alloy can withstand a stress of ~69 MPa (or 10,00 psi) for 100 hours without fracturing (1). Table VI lists some of the metals (best possible alloys) with respect to this criterion. The criterion does not have practical significance but it may be used to evaluate the effectiveness of an alloying element on pure metals. For magnesium, it can be observed that the best alloys (mainly rare earth containing and in the past thorium containing) can compete very well with aluminum alloys. Table 6. Best possible alloys of a base metal for elevated temperature service (ref 1)

base metal Mg Al Ti Fe (martensitic) Fe (austenitic) Ni

melting point (°C) 650 660 1704 1538 1538 1455

Temperature of useful strength for best alloy (°C)

% absolute melting point

343 288 649 732 871 1071

67 60 46 56 63 78

87 4.1

Yttrium-Containing Alloys

Low pressure and gravity casting alloys of magnesium have been traditionally used in aerospace applications. There are well-known creep-resistant alloy systems based on rare earth elements and yttrium such as WE43, WE54. These alloys do not contain aluminum and are grain-refined effectively with zirconium. Mg-Y-RE (Nd) alloys are amenable to precipitationhardening. The microstructures of WE54 and WE43 contain a eutectic phase at the grain boundaries in the as-cast state that has been characterized by Ahmed et al. (54) as β-phase, Mg14Nd2Y, with a face centered cubic crystal structure. Presence of another phase, Mg9Nd with a face-centered tetragonal structure after homogenization at 525°C, which was erroneously ascribed to the eutectic by Stulikova et al (55), was also reported by several other authors (56,57). Excellent room and elevated temperature tensile and creep properties of the alloy WE54 in the solution treated and aged condition are attributed (54, 58) to the formation of precipitates, namely β_, β_ (Mg12NdY), and β (Mg14Nd2Y). In addition to the standardized alloys in the Mg-Y-RE system, some experimental compositions have also been studied. Kamado et al (59) reported the ageing characteristics and tensile properties of some Mg-Gd-Y-Zr alloys. The increasing age hardening with decreasing Y/Gd ratio was again attributed to the formation of β″ and β′ precipitates. Another important example is the positive contribution of zinc to the creep resistance in experimental quaternary alloys in Mg-Y-Zn-Zr system as reported by Morgan and Mordike (60). An optimum amount of about 2%Zn was shown to result in better steady state creep rates than for the Mg-Y-Zr alloys without zinc. The best creep resistance in Mg-YZr alloys was obtained with 2%Zn and 2%Nd addition in the as-cast condition (60). 4.2

Alloys Containing Zinc and RE Addition

Possibly the most important gravity cast magnesium alloy systems for automotive applications to date have been the Mg-Zn-RE system (ZE and EZ alloys). An alloy belonging to this group is ZE41 that offers moderate strength up to 150°C after aging of the cast material. Better properties in terms of castability, freedom from microporosity, strength and creep resistance up to 250 °C (16) can be obtained with EZ33 (2.5-4.0% RE, 2.0-3.1%Zn, 0.51.0Zr). The grain boundaries in these alloys following aging are decorated with Mg-RE precipitates. It has been suggested that the eventual equilibrium precipitate is Mg12Ce or Mg27Ce2. The aging sequence is assumed to be similar to that in Mg-Nd alloys involving several stages including the formation of a coherent β"-phase that causes hardening (16). Since they do not contain aluminum, these alloys can be grain-refined by 0.6-0.7% zirconium addition. A study by Fisher (61) represents a very different approach to modification of precipitates in this alloy system, and has led to a successful alloy used in aircraft industry. Hydrogenation of ZE62 at 500 °C for 24 hours followed by 64hr of aging at 125 °C caused the intragranular precipitation of RE hydrides with a needle-like morphology resulting in remarkable properties (16, 61). 4.3

Alloys Containing Zinc and Copper

Mg-Zn-Cu alloys were developed for a combination of good ductility and elevated temperature performance (62, 63), where Cu imparts grain-refining effect and Zn good castability and ductility. The microstructure is two-phase with α-Mg and a truly lamellar eutectic with the composition β-Mg (Cu.Zn)2. Upon solution-treatment and ageing two

88 mainprecipitates are observed β1 needles and β2 plates. Addition of Mn prevents overaging at 250 °C and improves overall yield strength. Practical casting alloy developed from this system is ZC63 (6Zn-2.5Cu-Mn) with good creep strength, good castability but low corrosion resistance. The alloy has been evaluated in select automotive applications. 4.4

Silicon-Containing Alloys

The commercial Mg-Al-Si diecasting alloys are hypoeutectic alloys with respect to the Mg-Mg2Si pseudo-binary system. The eutectic point lies at 3.67% Mg2Si in pseudobinary Mg-Mg2Si. Any alloy to the hypereutectic side of this pseudo-binary system has not been considered commercially. It has been shown, however, that the hypereutectic alloys may be further strengthened by Al, Cu and RE additions via solid solution hardening (7). The use of Mg-Al-Si alloys in the sand and permanent-mold cast condition is difficult due to the existence of the Mg2Si in a coarse Chinese script morphology, which lowers ductility extensively. Pekguleryuz et al (8) have shown that the addition of Ca in the range of 0.080.14%Ca modifies the Mg2Si with resultant increase in ductility and tensile strength. (Fig.12).

(a)

(b)

Fig. 12. (i) Mg2Si Chinese-script morphology in AS41; (ii) 0.1% Ca modified AS41 with refined Mg2Si (8).

4.5

Recent Alloy Development in Magnesium Gravity Casting Alloys for Powertrain Applications

Two recent alloy development activities have the objective of developing cost-effective magnesium gravity-casting alloys with good creep resistance for automotive powertrain. One of these is the Mg-Zn-RE system (MEZ) alloy developed by MEL (64). The alloy has 2.5% RE and 0.5% RE. The permanent mold-cast microstructure, as characterized in a study by Bettles et al, comprises intermetallic particles of tetragonal Mg12(La0.43Ce0.57) at the grain boundaries (65). The alloys are age hardening with improved strength upon ageing. The second alloy development effort is from Dead Sea Magnesium and Volkswagen AG (48). The alloys are coded MRI 201S and MRI 202S. The alloy MRI 201S in the T-6 conditions with good mechanical properties and creep resistance and castability, is claimed to be a cost-effective substitution for WE43. The alloy MRI 202S with moderate strength and good creep performance is shown to be superior to ZE41 alloy. New gravity-casting alloy possibilities do exist with alkaline earth (Ca, Sr) additions as well. Very few gravity casting alloys have been explored in Mg-Al-(Alkaline Earth) and Mg-Zn-(Alkaline Earth) systems and these could provide a good starting point for new alloy development. In fact, without the constraints of diecastability, age-hardening and dispersion hardening can be effectively used to improve creep-resistance and elevatedtemperature performance.

89 4.6

Properties of Magnesium Gravity Casting Alloys

Again it is difficult to make a sound comparison of the various alloys from published data. A summary table is provided below (Table VII). USCAR program will be able to provide more rigorous comparative data on these alloys. Table 7. Properties of magnesium gravity casting alloys

Properties

WE43-T6 ZE41-T5 ZC63-T6 ASX411-F MEZ –T6 MRI 201S-T6 MRI 202S-T6

YS (MPa) 20 C 180 150C 175

140 120

158 134

75

76

170 170

150 145

260 210

220 170

240 180

175

107

260 245

250 220

6 7

5 22

4.5 14

7

2

6 11

7 15

160 60 0.10

50 20 3.1

63 *

-

-

160 75 0.10

100 40 0.12

UTS (MPa) 20 C 150 C Elongation (%) 20 C 150 C Stress to produce 0.2% creep-strain 200 C 250 C Corrosion 2 (mg/cm /day)

* Better than AZ91C

5

Synopsis

Magnesium diecasting alloys for elevated-temperature applications are coming of age. A number of aluminum containing alloy systems with single or combined additions of rareearth and/or alkaline-earth elements have been developed. Creep resistance in these alloys is attributed to grain boundary strengthening by intermetallic phases. Developing alloy compositions that avoid detrimental creep-induced precipitation which contributes to grain boundary migration is the key learning of the past decade. These new alloys are providing commercially viable alloys for automotive power-train applications such as oilpans, transmission-case and engine components. In the gravity and low-pressure casting area, the effort is still in its infancy and future R&D as well as interest on the part of car companies are expected to contribute to steady progress in this field. The role of RE’s and of yttrium in providing creep resistance is well known. New alternatives in Sr and Ca additions similar to diecasting versions can be explored in the future. Age-hardening, dispersion hardening can be effectively exploited for improving creep-resistance.

90

6

References

[1] R. E. Reed-Hill, Physical Metallurgy Principles, 2nd Edition, D. Van Nostrand, New York, 1973. [2] M. F. Ashby, D. R. H. Jones, Engineering Materials 1, Pergamon Press, Oxford, 1980. [3] E. F. Emley, Principles of Magnesium Technology, Pergamon Press, London, 1966. [4] G.V. Raynor, The Physical Metallurgy of Magnesium and its Alloys, Pergamon Press, New York, 1959 [5] M.S. Dargusch, G.L. Dunlop and K. Pettersen, Magnesium Alloys and Their Applications, Proc. Volume sponsored by Volkswagen AG, B.L. Mordike and K.U. Kainer, Eds., Werkstoff-Informationsgesellschaft, Frankfurt, Germany, 1998, pp. 277-282. [6] Y. Terada, N. Ishimatsu, R. Sota, T. Sato, K. Ohori, “Creep Characteristics of CaAdded Die-Cast AM50 Magnesium Alloys,” Magnesium Alloys 2003, Materials Science Forum, Trans Tech, Switzerland, vols. 419-422, 2003, pp. 459-464. [7] S. Beer, G. Frommeyer and E. Schmid, Magnesium Alloys and Their Applications, DGM Conf. Proc., B.L. Mordike and F. Hehmann, Eds., April 1992, pp. 317-324. [8] M.Pekguleryuz, A. Luo, C. Aliravci, “Modification of Sand Cast AS41 Magnesium Alloy,” Proc. Int. Symp. Light Metals Processing and Applications, The Metallurgical Society of CIM, Quebec City, Canada, Aug-Sep 1993, pp. 409-416. [9] G.S. Foerster, Designing Diecasting Alloys, Light Metal Age, v.30, 9-10, 1972, pp. 11-13. [10] M.S. Dargush and G.L. Dunlop, Proc. of Materials 98, Institute of Materials Engineering, Australasia Ltd., ed. Michael Ferry, Wollongong, Australia, 1998, p.579. [11] W. Blum, B. Watzinger and P. Weidinger, Magnesium Alloys and Their Applications. B.L. Mordike and K.U. Kainer, Eds., Werkstoff-Informationsgesellschaft, Frankfurt, Germany, 1998, p. 49-60. [12] T. Kr. Aune and T.J. Ruden, SAE 920070, Detroit, Michigan, February, 1992. [13] K. Pettersen, H. Westengen, J.I. Skar, M. Videm, L-Y. Wei, “Creep Resistant Mg alloy Development,” Magnesium Alloys and Their Applications. K.U. Kainer, Ed., DGM, Wiley-VCH, Weinheim, Germany, 2000, pp. 29-34. [14] T.Kr. Aune and H. Westengen, SAE Technical Paper No 950424, Detroit, Michigan, February, 1995. [15] William Mercer II, SAE Technical Paper No 900788, Detroit, Michigan, February, 1990. rd [16] I.J. Polmear, Light Alloys, Metallurgy of Light Metals, 3 addition, Arnold, London, 1995. [17] O. Holta, M. Videm, H. Westengen and D.L. Albright, Diecasting Innovation, by North American Diecasting Association (NADCA), Indianapolis T95-053, Indiana, October, 1995, p.177-182. [18] G. Pettersen, H. Westengen, R. Hoier and O.Lohne, Mat. Sci. and Eng., A207, 1996 pp. 115-120. [19] L.Y. Wei, and G.L. Dunlop, J. Alloys and Compounds, 232, 1996, p. 264. [20] B. R. Powell, V. Rezhets, M.P. Balogh, and R.A. Waldo, “Microstructure and Creep Behavior in AE42 Magnesium Die-Casting Alloy”, Journal of Metals, TMS, August 2002, pp. 34-38.

91 [21] A.A. Luo and M.O. Pekguleryuz, “Cast Magnesium Alloys for Elevated Temperature Applications,” J. Mat. Sci., vol. 29, 1994, pp. 5259-5271. [22] P. Bakke, K. Pettersen, H. Westengen, “Enhanced Ductility and Strength Through RE Addition to Magnesium Diecasting Alloys,” 2003 Magnesium Technology, TMS, San Diego, March 2003, pp. 171-176. [23] F. Hollfigl-Rosta, E. Just, J. Kohler and H.J. Melzer, Light Metals Age, August 1980, p.22-29. [24] H.J. Fuch, “Magnesium Alloys Having a High Resistance to Permanent Creep Deformation at Elevated Temperatures”, British Patent 847, 992, 1960. [25] M.O. Pekgüleryüz and A. Luo, “Creep Resistant Magnesium Alloys for Diecasting”, PCT (Patent Cooperative Treaty) application No. PCT/CA96/00091 published August 22, 1996. [26] M. Pekguleryuz, A. Luo, "Creep Resistant Magnesium Alloy”. Patent WO96/25529 (1995) [27] R. Ninomiya, T. Ojiro and K. Kubota, Acta Metall., vol. 43, No.2, 1995, pp.669-674. [28] M. O. Pekguleryuz and J. Renaud, “Creep Resistance in Mg-Al-Ca Alloys,” 2000 Magnesium Technology, H. Kaplan, J. Hryn and B. Clow, Eds. TMS, 2000, pp. 279-84. [29] H. Gjestland, G. Nussbaum and G. Regazzoni, Mat. Sci. And Eng., A134, 1991, pp. 1197-1200. th [30] F. Holbrigl-Rosta, E. Just, J. Kohler and H.J. Melzer, Proc. 37 Annual World Magnesium Conf., IMA, New York, 1980, p. 38. [31] A. Luo and T. Shinoda, “Magnesium Alloys Having Superior Elevated-Temperature Properties and Die Castability,” US Patent 5,855, 997, January 5, 1999. [32] EP 0799901A1, Mazda 1997. [33] A. A. Luo, “Recent Magnesium Alloy Development for Automotive Powertrain Applications,” Magnesium Alloys 2003, Materials Science Forum, Trans Tech, Switzerland, vols. 419-422, 2003, pp. 57-65. [34] M. Pekguleryuz, P. Labelle, “Magnesium-Based Casting Alloys Having Improved Elevated Temperature Performance,” US Pat. No. 6,322,644, Nov 27th, 2001 and International Patent PCT WO 01/44529 [35] M. O. Pekguleryuz, “Development of Creep Resistant Magnesium Diecasting Alloys,” Magnesium Alloys 2000, Materials Science Forum, Y. Kojima, T. Aizawa, S. Kamado, Eds. Trans. Tech. Publication, Switzerland, 2000, pp. 131-130. [36] M. O. Pekguleryuz, E. Baril, “Development of Creep Resistant Magnesium-AluminumStrontium Alloys, 2001 Magnesium Technology, J. Hryn, Ed., TMS, New Orleans, March 2001 pp.119-136. [37] P. Labelle, M. Pekguleryuz, M. Lefebvre, R. Bouchard, “New Aspects of Temperature Behavior of AJ52x, Creep Resistant Magnesium Alloy,” SAE Technical Paper 02M-161, 2002. [38] D. Argo, M. O. Pekguleryuz, P. Labelle M. Dierks, T. Sparks, T. Waltematte, “Diecastability and Properties of Mg-Al-Sr Based Alloys,” 2001 Magnesium Technology, J. Hryn, Ed., TMS, New Orleans, March 2001 pp.131-125 [39] M. Pekguleryuz, P. Labelle, D. Argo, M. Dierks, T. Sparks, T. Waltematte, “CreepResistant Magnesium Diecasting Alloys” Proc. 57th Ann. World Magnesium Conf., IMA, Brussels, May 2001, pp. 27-33.

92 [40] D. Argo, M. Pekguleryuz, P. Labelle, P. Vermette, R. Bouchard and M. Lefebvre “Process Parameters And Diecasting of Noranda’s AJ52 High Temperature Mg-AlSr Alloy,” 2002 Magnesium Technology, H. Kaplan, Ed., TMS, Seattle, Feb 2002, pp. 87-93. [41] Yemi Fasoyinu, Terri Castles, Réal Bouchard, Mahi Sahoo, M. Pekguleryuz, and Pierre Labelle, “Die cast Magnesium Alloys AE42 and AJ52x for High Temperature Applications,” 2003 Magnesium Technology, TMS, San Diego, March 2003, p. 207. [42] M. Pekguleryuz, P. Labelle, E.Baril, D. Argo, ‘’Magnesium Diecasting Alloy AJ62x with Superior Creep Resistance and Castability,” 2003 Magnesium Technology, TMS, San Diego, March 2003, pp. 201-207. [43] M. Pekguleryuz, P. Labelle, E.Baril, D. Argo, “’Magnesium Diecasting Alloy AJ62x with Superior Creep Resistance, Ductility and Castability,” SAE Technical Paper 2003-01-190, Detroit, March 2003. [44] K. Samato, Y. Yamamoto, N. Sakate, and S. Hirabara, “Heat-Resistant Magnesium Alloy Member,” EP 0 799 901 A1 August 10 1997. [45] S. Koike, K. Wasizu, S. Tanaka, T. Baba and K. Kikawa, SAE Technical Paper 20001-1117, Detroit, March, 2000. [46] M. Pekguleryuz, “Creep Resistant Magnesium Diecasting Alloys Based on Mg-Al(Alkaline Earth Element) Systems", Materials Transactions, Special issue on Platform Science and Technology for Advanced Magnesium Alloys, Japan Inst. Metals, 2001, Vol. 42, No. 7, pp. 1258-1267. [47] B. Bronfin, E. Aghion, S. Schuman, P. Bohling, and K.U. Kainer, “Magnesium Alloy for High Temperature Applications,” US Patent 6,139,651, Oct 31, 2000. [48] E. Aghion, B. Bronfin, H.Friedrich, S. Schumann and F. von Buch. New Magnesium Alloys For High Temperature Applications, 2003 Magnesium Technology, TMS, San Diego, March 2003, pp. 177-182. [49] M. Lefebvre, M. Pekguleryuz, P. Labelle, “Magnesium-Based Casting Alloys Having Improved Elevated Temperature Performance,” US Pat. No 6, 342,180, Jan 29, 2002 and International Patent PCT WO 02/099147 [50] B.R. Powell, V. Rezhets, A.A.Luo, J.J. Bommarito, and B.L. Tiwari, “Creep Resistant Magnesium Alloy Die Casting, “ US patent 6,264,763, July 24, 2001. [51] A. Luo, M. Balough, and B. R. Powell, SAE Technical Paper 2001-01-0423, Detroit, Michigan, 2001. [52] A. A. Luo, “Recent Magnesium Alloy Development for Automotive Powertrain Applications,” Magnesium Alloys 2003, Materials Science Forum, Trans Tech, Switzerland, vols. 419-422, pp. 57-65. [53] A. P. Druschitz, E. R. Showalter, J.B. McNeill, D. L. White, “Evaluation of Structural and High-temperature Magnesium Alloys”, SAE Technical Paper 2002-010080, Detroit , MI, Feb 2002 [54] M. Ahmed, G.W. Lorimer, P. Lyon, and R. Pilkington, Magnesium Alloys and Their Applications, B.L. Mordike & F. Hehmann, Eds., DGM Informationsgesellschaft, Verlag, April 1992, p.301-308. th [55] I. Stulikova, P. Vostry, B. Smola and B.L. Mordike, Proc. of ‘5 Int. Conf. on Heat Treatment of Materials, IFHT, Budapest, 1986, p.1993. [56] J.M. Worrall, M.Sc. Thesis, Manchester University, 1986 [57] P. Vostry, I. Stulikova, B. Smola, M. Cieslar and B.L. Mordike, Z. Metallkunde, 1985, p. 340

93 [58] M. Ahmed, G.W. Lorimer, P.Lyon, and R. Pilkington, Magnesium Alloys and Their Applications, B.L. Mordike & F. Hehmann, Ed., DGM Informationsgesellschaft, Verlag, April 1992, p.251-257. [59] S. Kamado, Y. Kojima, S. Taniike and S. Hama, ibid, pp.169-174. [60] J.E. Morgan and B.L. Mordike, Eds., Met. Trans., vol. 12A, 1981, pp.1581-1585. [61] P.A. Fisher et al., Foundry, 95(8), 68, 1967 [62] W. Unsworth, “A New Magnesium Alloy for Automobile Applications,” magnesium in the Auto Industry: prospects for the Future, Proc. 44th An. World Magnesium Conf. IMA, JLMA, May 1987, pp. 22-27 [63] W. Unsworth, ‘New ZCM Magnesium Alloy,” SAE technical Paper 880512, 1988. [64] J. F. King, P. Lyon, and K. Nuttall, US Patent 6,193,817, 2001. [65] C.J. Bettles, K. Venkatesan and J.F. Nie, “Microstructure and Mechanical Properties of MEZ Casting Alloy,” Magnesium Alloys 2003, Materials Science Forum, Trans Tech, Switzerland, vols. 419-422, 2003, pp. 273-277.

94

Russian Ultralight Constructional Mg–Li Alloys. Their Structure, Properties, Manufacturing, Applications Fedor M. Elkin, Valentin G. Davydov (All-Russia Institute of Light Alloys, Moscow)

Peculiarities of Mg–Li alloys The Mg–Li alloys are the ultralightweight constructional metallic materials having the 1.35–1.65 g · cm-3 density, that 1.5–2.0 times less than that of aluminium alloys. These density meanings are similar to constructional plastic density. The second principal advantage of the Mg–Li alloys is the fact that they have high ductility. It is caused by the isotropic body-centered cubic (bcc) lattice of them instead of hardly deformed at ambient temperatures and anisotropy hexagonal close-packed (hcp) lattice of traditional magnesium alloys.

Commercial Mg-Li alloys In Russia (USSR) the original multicomponent compositions of wrought and cast Mg–Li alloys were developed. Two wrought Mg–Li alloys were standardized in 1983 (Russian Standard “Magnesium Alloys” — GOST 14957). These alloys have of marks MA18 and MA21 (ИМВ2). The alloy MA21 was developed on base of (hcp + bcc) lattices and that MA18 was developed on the base of bcc lattice. Table 1. Chemical composition limits of MA21 and MA18 wrought Mg–Li alloys.

Alloy Composition, wt.% (Mg – remainder) Li

Al

Zn

Cd

Mn

Ce

Na max

K max

Fe max

MA21 7.5–9.0 4.3–5.3 1.0–2.0 4.0–5.0 0.03–0.1 0.01–0.15 0.005 0.005 0.01 MA18 10–11.5 0.5–1.0 2.0–2.5 — 0.1–0.4 0.15–0.35 0.005 0.005 0.01 The commercial Mg–Al alloys MA18 and MA21 (ИМВ2) were developed in 60–70’th years of the last century by three leading Soviet Union scientific institutes, engaged with researches of light alloys — VILS (All Russia Institute of Light Alloys), IMET (Baikov Institute of Metallurgy) and VIAM (All-Russia Institute of Aviation Materials). The fundamental results of these investigations of Mg–Li alloys were published in Soviet Union in two books [1–2] and two reference books [3–4]. The main feature of the compositions of the Mg–Li-Al alloys with sufficiently high lithium contents (5–15 wt.%) is non-stability of their mechanical properties during expo-

95 sure even at ambient temperatures. Such the Mg–Li–Al alloys are amenable to hardening during ageing although they overage and soften at relatively low temperatures (e.g. 5070 °C) as a consequence of the abnormally high mobility of lithium atoms and vacancies. Somewhat greater stability and creep resistance of the Mg–Li–Al alloys have been achieved by adding alloying elements such as Zn and Cd, which have sufficiently high solid solubility in the phases based on both magnesium (α-phase) and lithium (β-phase). The small amount of Mn and Ce alloying elements show the positive influence on the corrosion resistance and on the technological properties at continuous casting. The main physic-mechanical properties of wrought semi-products from MA21 and MA18 alloys are presented in table 2. Table 2. The main physic-mechanical properties of MA21 and MA18 wrought Mg–Li alloys.

Alloy

Crystal lattice

20 Density, σB, MPa σ0,2, MPa σ-0.2, MPa δ, % E, GPa σ 0.2 /1000 , MPa -3 g ⋅ cm

MA21 hcp+bcc 1.60 (α+β) phases MA18 bcc 1.48 (β phase)

200–260 130–220 150–220

6–20 45–46 80

140–210 90–180

10– 45–46 60 30

100–180

The minimum strength properties in the table 2 concern to the big forgings, die forgings and extruded bars and large section shapes. The maximum strength properties in the table 2 concern to the warm rolled thin sheets and hot extruded thin wall shapes and tubes. The MA18 alloy microstructure consists of the β-phase grains with the fine-disperse (Al, Zn)Li-phase inclusions. Besides the Mn and Ce phase containing particles are (Fig. 1).

a

b

c

d

Figure 1. Microstructure of the MA18 alloy: a — billet, ×100, b — billet, × 7000, c — 2 mm sheet, × 340, d — thin-wall extruded profile, × 340

96 The microstructure of the MA21 alloy is more complex than that of MA18 alloy. It consists of the two main phases (α + β) in approximately equal quantities with the phase (Al, Zn, Cd) Li fine-disperse metallic particles. The grain size of α and β phase particles depends very much on the crystallization rate. Fig. 2 shows microstructure of section through Electron Beam Welded (EBW) in MA21 die forging (× 250).

I

II

III

Figure 2. Microstructure (× 250) of EBW weld junction in MA21 die forging: I — die forging, II — transitional zone, III — weld metal.

As Fig. 2 shows the fine grains of α and β phases are formed in weld metal where more rapid crystallization takes place. The possibility to reduce fragments of the (α + β) grains in MA21 alloy, e.g. by granular metallurgy, permitted to increase strength of wrought semi-products by 1.2 – 1.5 times as compared with traditional technology of production using billets. The ultralightweight MA18 alloy (γ = 1.48 g ⋅ cm-3) and LA141A (USA alloy) with γ = 1.35 g ⋅ cm-3 have low strength and creep resistance. Therefore these alloys have been intended for special applications such as components for spacecraft, housings and cases of computers, notebooks, mobile telephones, electron equipment, bogy armor etc. These alloys have useful special properties as high heat capacity, high impact strength, high stiffness, high armor properties, good mechanical properties at cryogenic temperatures etc. The wrought weldable Mg–Li alloy MA21 (ИМВ2) has no similar analogue in other countries. This alloy has the excellent combination of low density (γ = 1.6 g ⋅ cm-3), sufficient high strength, good ductility and satisfactory corrosion resistance. The yield strength of big forgings and die forgings (weight about from some hundred kg up to one ton) from MA21 alloy (σ0.2 = 130–160 MPa) is not less that the same characteristic of the high strength weldable Russian aluminium АМГ6 alloy (type of AA5056 alloy of USA) with the density γ = 2.65 g ⋅ cm-3. The preference of MA21 alloy as compared with traditional weldable without heat treatment magnesium (MA2/AZ31B) and aluminium (ΑΜΓ6/AA5056) alloys is more significant considering specific yield strength and modulus (stiffness) of big forgings, die forgings and extruded large section shapes (Table 3). The alloys MA18 and MA21 have good weldability performed by many methods of weld, including EBW, up to thickness 40–60 mm. The big wrought semi-products of MA18 and MA21 alloys do not have anisotropy of mechanical properties, on contrast with traditional magnesium alloys.

97 The wrought semi-products (sheets, plates, extruded shapes, forgings and die forgings) of the MA21 and MA18 alloys were used in Defence and Aerospace industry in such welded constructions where the considerable weight saving is of great importance. Table 3. The density, specific yield strength and stiffness of weldable without heat treatment standard Mg–Li, Mg and Al alloys. (The big forging, die forging and extruded large section shapes)

Alloy Russia/USA MA21 MA2/AZ31B АМГ6/AA5056

γ, g⋅cm-3 1.60 1.78 2.65

σ0.2/γ 8.8 6.2 5.2

σ-0.2/γ 10.0 4.0 5.2

E/γ 3000 2300 2600

Tentative Mg-Li alloys At present in the world the ultraweightlight wrought Mg–Li alloys with sufficient strength and corrosion resistance are interesting for civil branches of industry e.g. automobile and electronic industries — wheels and components of cars interior, housings and cases of computers, notebooks, mobile telephones, frames video-photo cameras etc. In Russia VILS and IMET developed a number of new tentative Mg–Li alloys having up to 8 wt. % of lithium on the base of (α + β) phases with sufficient strength and fatigue resistance for wheels of sports cars of Formula I [5]. Table 4. Nominal compositions and typical room temperature physic-mechanical properties of tentative Mg–Li and standard Mg–Li and Mg alloys.

Crystal Density, σв, σ0.2, lattice g ⋅ cm-3 MPa MPa (phase) ИМВ1 Mg-5.5Li-5.5Al-1Zn- hcp (α) 1.65 270 210 0.5Mn-0.8Sn ИМВ12 Mg-6Li-Y, Mn, Zn hcp+bcc 1.57 220 170 (α+β) ИМВ14 Mg-6.5Li-Al, Zn, Cd, – “ – 1.58 250 170 Mn MA21 Mg-8Li-5Al-4Cd–“– 1.60 240 180 1.5Zn-0.1Ce МА18 Mg-11Li-2.2Znbcc (β) 1.48 170 120 0.7Al-0.25Mn-0.25Ce LA141A Mg-14Li-1Al-0.15Mn – “ – 1.35 150 110 * Mg-15Li-5Zn-0.5Al- – “ – 1.38 180 – REM,Y *МЛ5пч Mg-8.5Al-0.5Znhcp 1.81 180 100 (AZ91D) 0.35Mn (Mg) Mg-4.5Al-1.2Zn–“– 1.79 280 200 МA21пч 0.4Mn AZ31B Mg-3.5Al-0.5Zn–“– 1.78 260 180 0.35Mn Alloy

Composition, wt. %

Footnotes: * cast alloys for high pressure die castings.

σ-0.2, δ, aн, HB E, MPa % kj ⋅ m GPa 2

230

5

180

60

86

44

15 150

72

45

180

12 250

75

45

190

12 120

75

45

130

20 250

65

45

120 –

20 300 5 100

60 70

43 –

–

2

50

63

42

120

10 80

82

42

110

12 90

80

42

98 Table 4 shows the main characteristics and the nominal compositions of tentative and standard Mg–Li alloys and other standard Mg alloys for the possible application in the automobile and electronic industries. In the end of last century VILS made wheels for sport cars of Formula I from die forgings of tentative Mg–Li alloys (ИМВ12 and ИМВ14) and standard Mg–Li alloy (MA21). The strength of cast Mg–Li alloys is rather similar the strength of wrought Mg–Li alloys of the same compositions, it is their peculiarity. The scientists of Perm Polytechnical Institute (Ural) developed tentative Mg–Li alloys for high pressure die casting [6–7]. These alloys have chemical compositions: Mg–12÷16% Li–3÷7 % Zn–0.5÷2 % Al–0.5÷2 % R.E.M and Y. The die cast samples of these alloys showed following properties: γ = 1.3–1.45 g⋅cm-3, σв = 170–220 MPa, δ = 2–5%, HB = 65–75. These Mg–Li alloys were characterized by good castability and did not showed the hot cracks during casting. At present in Russia commercial cast Mg–Li alloys does not use.

Mg-Li alloys technology Constructional Mg–Li alloys contain high lithium (up to 15 wt.%) which is chemical active and toxic alkaline element. This causes necessity to create of the special manufacture technologies. These industrial technologies include fluxless melting, refining, continuous casting, machining, hot, warm and cold working, heat treatment, recycling and other processes. VILS have developed these technologies and has industrial equipment for wide range semiproducts from wrought Mg–Li alloys production. This equipment includes melting–casting aggregates with induction crucible furnaces of 0.5, 1.7 and 3.0 tons capacity for casting billets having 370–680 mm diameters and 165 × 650 mm sections slabs with length and weight up to 5 m and up to 3 tons, respectively. The scheme and photo of Mg–Li aggregate with 1.7 ton capacity melting crucible furnace are shown in Fig. 3. shielding gas shielding gas

a

b

Figure 3. The scheme (a) and photo (b) of Mg–Li aggregate with 1.7 ton capacity furnace. 1 – Induction crucible melting furnace. 2 – Electromagnetic pump. 3 – Continuous casting equipment. 4 – Control equipment for shielding gas.

99 From MA21 alloy 680 mm in diameter billets were made forgings and die forgings with diameters up to 2500 mm and up to one ton by weight. Mg–Li alloys have excellent formability during hot and warm working. In this aspect they are better than traditional magnesium alloys. The temperatures of hot working Mg–Li alloys beneath 320 °C because higher temperatures would result in inflammation.

References [1] M. E. Dritz, F. M. Elkin. I. I. Guriev, B. I. Bondarev et al. Magnesium – Lithium Alloys. Russia, Moscow, Metallurgy. 1980. P. 139. [2] M. E. Dritz, Z. A. Sviderskaia, F. M. Elkin, V. F. Trokhova. Ultralight Constructional Alloys. Russia, Moscow, Nauka, 1972. P. 144. [3] F. M. Elkin. Magnesium and Magnesium Alloys in Reference Book “International Translator of Modern Alloys” (Ed: V. J. Kershenbaum) International Academy of Engineering. Russia, Moscow, 1993, Volume III. P. 453–532. [4] F. M. Elkin. Ultralight Alloys in Reference Book “Magnesium Alloys” (Ed: M. B. Altman et al.) Russia, Moscow, Metallurgy. 1978. Volume I. P. 232. [5] Elkin F. M., Naryshkina N. A., Bondarev B. I. Et al. Russian Patent № 2048578, 1995. [6] L. V. Nikulin, S. P. Novikov et al. in “Litie, cristallizacia magnievyh i aluminievyh splavov”, Russia, Perm, 1980. P. 22–27. [7] M. A. Tomsinskaia, S. Ja. Popov in “Litie, cristallizacia magnievyh i aluminievyh Splavov”, Russia, Perm, 1980. P. 28–31.

100

The Effect of Grain Size on the Mechanical Properties of AM-SC1 G. Dunlop1, C.J. Bettles2, J.R. Griffiths2, K. Venkatesan2, Lihui Zheng3 and Ma Qian3 1

Australian Magnesium Corporation, Milton, Australia. CRC for Cast Metals Manufacturing (CAST), CSIRO Manufacturing & Infrastructure Technology, Melbourne, Australia. 3 CRC for Cast Metals Manufacturing (CAST), The University of Queensland, St Lucia, Australia 2

Introduction Australian Magnesium Corporation, in collaboration with the Cooperative Research Centre for Cast Metals Manufacturing (CAST) and Magnesium Elektron Limited, has developed a magnesium alloy, AM-SC1, which has been specifically designed for engine block applications[1]. This alloy has been used for the engine block of the Genois LE turbo charged diesel injection motor developed by AVL List [2]. AM-SC1 was developed for sand or permanent mould casting processes, and the mechanical properties may be optimised by the use of an appropriate heat treatment. However, the grain size achieved during casting also significantly affects both the strength and creep resistance. To ensure that a uniform grain size is achieved, irrespective of the thickness of the cast section, a grain refining agent is included in the alloy composition. The level of addition determines the grain size and, ultimately, the mechanical properties of the casting. The strengthening effect of grain boundaries at room temperature is generally described by the Hall-Petch relationship [3, 4]. σf = σo + k d

-1/2

(1)

Similarly, the steady state creep rate can be described by p

n

Q

⎛ b ⎞ ⎛σ ⎞ − ε& = A ⎜ ⎟ ⎜ ⎟ e RT ⎝d⎠ ⎝G⎠

(2)

which predicts that the creep rate can be reduced by increasing the grain size. It can be seen that a high yield strength requires a small grain size while creep resistance requires a large grain size. The aim of this investigation was to determine the extent of grain size effects on the offset yield strength (tensile and compressive), tensile creep, bolt load retention behaviour and fatigue strength of AM-SC1.

Experimental Methods Four plates of AM-SC1, 160 mm x 180 mm x 30 mm (l x b x d) were sand cast, each with a different level of Zr addition (including zero addition) using AM-cast grain refiner [3],

101 to achieve grain sizes ranging from 110 μm to 536 μm. A fifth plate, also containing AMcast, was cast using a chill to achieve the smallest grain size of 46μm. It should be noted that the grain refiner does not, in itself, contribute to the mechanical properties of the material. All five plates were given the recommended T6 heat treatment for the alloy. The grain size was determined by a linear intercept method (ASTM E112-96), using three concentric circles of total circumferential length 900 mm. The polished surfaces were etched with a picric acid-acetic acid etchant to enhance the definition of the individual grains under polarised light. Tensile and compressive yield strengths(0.2% offset) were determined at 293, 373, 423 and 450 K, according to ASTM E8 and E9 respectively. Tensile creep testing was carried out under constant load conditions, at 423 and 450 K and 90 MPa. The bolt load retention behaviour was determined using the method described by Pettersen and Fairchild [6] at 423 and 450 K and loads of 8 kN and 11 kN. The results presented are the average of at least duplicate tests. The fatigue strength was determined at R = -1 at room temperature 6 only. The fatigue strength at 10 cycles was determined by the staircase method [7] using at least 15 specimens; the strength at 107 cycles was also estimated by the staircase method but using fewer specimens; the strength at 105 cycles was estimated by interpolation, typically using at least 5 data points. The specimen geometries for all the above tests are described elsewhere [8].

Results and Discussion Microstructure and Grain Size The as-cast microstructures of the five alloys are shown in Figure 1. The microstructure of the plate containing no grain refining agent is cellular in nature, with a significant amount of an interdendritic intermetallic phase (Figure 1(a)). The mean grain size was 536 μm, with a standard deviation of 156 μm (see Table 1), and grains in excess of 1mm were not uncommon. As the grain refining level was increased, the cellular structure became more regular (Figure 1 (b) and (c)), and ultimately took on the conventional appearance of a fully grain refined alloy (Figure 1(e)). At the same time the variability of the grain size decreased.

d

e

c

100μm

b

200μm

a

Figure 1. Optical micrographs of the five plate castings showing the differences in grain morphology and size. Note that figures 1(a)–(d) (Alloys 1-4) are at the same magnification and that 1(e) (alloy 5) is at a higher magnification.

102 Table 1. Grain sizes of the five AM-SC1 plates (where Alloy 1 has no grain refinement and Alloy 5 has the maximum addition plus a chill).

Grain size (μm) 536 ± 156 337 ± 96 235 ± 28 110 ± 9 46 ± 3

Alloy 1 Alloy 2 Alloy 3 Alloy 4 Alloy 5

Tensile and Compressive Yield Strengths (0.2% offset) Plots of the tensile and compressive yield strengths, σ0.2%T and σ0.2%C, as functions of grain size at various temperatures are shown in. These plots show that, in both tension and compression, the yield strength - grain size relationship can be represented by an equation of the form σ0.2% = A – Kd

(3)

tensile yield strength (MPa)

140

compressive yield strength (MPa)

where d is the grain diameter listed in Table 1. This relationship is in marked contrast to the more usual d-1/2 one of Equation 1(see Figure 2(c) for the Hall-Petch plot, demonstrating the deviation from linearity). Also shown in are the specifications for the tensile yield strength at room temperature and 450 K (horizontal dotted lines) required by the engine designer [1]. It is clear that, to achieve this strength specification, the grain size needs to be less than 150 μm. (a)

130

293K 450K

120 110 100 90

293K

80

373K

70

423K 450K

60 0

100

200

300

400

500

600

150

(b)

140 130 120 110 100 90

293K

80

373K 423K

70

450K

60 0

200

400

600

grain size (µ m )

grain size (µm ) 140

yield stress (MPa)

130

(c)

Figure 2. The effect of grain size on (a) the tensile yield strength and (b) the compressive yield strength at room and elevated temperatures. Horizontal dashed lines indicate tensile yield strength design requirements at 293 K and 450 K for the development of an engine block for a turbo charged direct injection diesel engine Also shown in (c) is the standard Hall-Petch plot of the data

120 110 100

tensile compressive

90 80

0.04

0.06

0.08

d

0.10 -0.5

-0.5

(μm )

0.12

0.14

0.16

103 The effect of temperature over the range 293 – 450 K is more pronounced in tension than in compression. Indeed the compressive yield strength is essentially independent of temperature over the range tested. It should also be noticed that the compressive values are similar to the room temperature tensile values, indicating that the alloy displays isotropic behaviour at room temperature irrespective of the grain size or morphology. At elevated temperatures the alloy is stronger in compression than in tension. The tensile yield strength results at the three elevated temperatures are very similar, with the exception of the data point for the 235 μm material at 423 K. The error bar on this particular data point is large and it is likely that further testing would remove the anomaly. The failure of the Hall-Petch relationship to describe the yield strength is remarkable and reasons for it are subject to our continuing research. Tensile Creep Behaviour To determine the role of grain size, constant load creep tests were carried out at 90 MPa at two temperatures: 423 and 450 K. The steady state creep rates were determined over the time interval 400-500 hours (Table 2). At 450 K the three largest grain sized alloys all showed increasing creep rates as the grain size decreased (Figure 3), as predicted by Equation 2. However as the grain size was reduced further (i.e. smaller than 200 μm) the creep rate decreased. Similar effects are found at 423 K with creep rates an order of magnitude less. Table 2. The steady state creep rates for AMSC1 with various grain sizes.

535 337 235 110 46

-11

1.7 x 10 2.8 x 10-11 5.1 x 10-11 4.2 x 10-11 2.0 x 10-11

-10

3.3 x 10 5.3 x 10-10 5.5 x 10-10 4.4 x 10-10 1.5 x 10-10

steady state creep rate (/s)

423K 90MPa 450K 90MPa Grain Size Steady State Steady State (μm) creep rate (s-1) creep rate (s-1)

6.00E-10 5.00E-10 4.00E-10 3.00E-10 2.00E-10 1.00E-10 0.00E+00

0 effect of 200grain size400 600 state creep Figure 3. The on the steady grain siz rates of AM-SC1 at 450K ande (microns) 90 MPa.

Bolt Load Retention Effects In terms of the overall behaviour of a fastened joint in an engine, the critical loads are PI, the initial load applied at room temperature, and PF, the load once the test assembly has returned to room temperature. The amount of relaxation at the end of a heating and cooling cycle, PF/PI, is the most important factor. Results are shown in Figure 3 for the two conditions 423 K 11 kN and 450 K 11 kN (a more detailed description of the effect of grain size and load on the BLR behaviour can be found in another paper presented at this conference [9]). At 423 K, the bolt load retention behaviour is relatively independent of grain size. The somewhat low PF/PI value at a grain size of 235 μm is an artefact since the initial applied

104

% Retained Load (P F/PI)

load was higher than the nominal 11 kN. At 450 K, the bolt load retention behaviour is relatively constant for the three smallest grain sizes. At larger grain sizes there is a drop in the load retained. This is probably because, as the grain size increases, the compressive yield strength decreases and, under the conditions of this test, the maximum stress approached the yield strength at a grain size of 337 μm and indeed exceeded it at 536 μm. 100 90 80 70 60 50 40 30 20 10 0 46

110

235

337

536

grain size (m icrons) 450K

423K

Figure 3. The effect of grain size on the retained load at room temperature after testing at 423 and 450 K and a load of 11 kN.

Fatigue The results from the staircase tests for lives of 106 and 107 cycles and for the interpolation for a life of 105 cycles are given in and are plotted in Figure 4. The data are for a probability of failure of 0.5 and the ± numbers are for failure probabilities of 0.1 and 0.9. As with the yield strength data the fatigue strengths decrease approximately linearly with increasing grain size. To test whether the fatigue strengths are, in fact, determined by the tensile properties the data are normalised with respect to the tensile yield strengths and ultimate tensile strengths and summarised in Table 4. It can be seen that there is good correlation with both yield strength and a slightly better one with ultimate tensile strength. 5

6

Grain size, Mean fatigue strengths, MPa, at various lives μm cycles 536 337 235 110 46

105

106

107

62 71 79 90 93

45 ± 15 56 ± 18 64 ± 15 72 ± 21 76 ± 16

37 ± 8 53 ± 11 56 ± 8 67 ± 14 ND

7

Strength, MPa

Table 3. Fatigue strengths at 10 , 10 and 10 cycles

100 95 90 85 80 75 70 65 60 55 50 45 40 35 30

105

106 107

0

100

200

300

400

500

600

grain size, microns

Figure 4. Fatigue strengths at various lives plotted against grain size.

105 Table 4. Fatigue strengths normalised with respect to the tensile yield strength and ultimate tensile strength.

Grain size, (Fatigue strength)/(Yield strength) μm at various lives 105 cycles 106 cycles 107 cycles 0.72 0.52 0.43 536 337 0.66 0.51 0.50 235 0.68 0.54 0.48 110 0.71 0.58 0.53 46 0.72 0.58 ND* mean ± s.d 0.70 ± 0.03 0.55 ± 0.03 0.49 ± 0.04

(Fatigue strength)/(Ultimate tensile strength) at various lives 105 cycles 106 cycles 107 cycles 0.54 0.39 0.32 0.48 0.37 0.36 0.48 0.38 0.34 0.44 0.36 0.33 0.45 0.36 ND* 0.48 ± 0.04 0.37 ± 0.01 0.34 ± 0.02

*ND not determined

Conclusions The dependence upon grain size of tensile and compressive yield strengths and creep strength of the high temperature sand casting alloy AM-SC1 does not obey the Hall-Petch relation but, instead follows a linear relationship with grain size of the type: σ0.2% = A – Kd. Secondary creep rates reach a maximum value at an intermediate grain size. Bolt load retention is relatively independent of grain size at 423 K in the range from 46 μm to 536 μm. At 450 K the BLR was found to decrease for grain sizes above 235 μm. The room temperature fatigue strength decreased with increasing grain size and there was a strong correlation between fatigue strength and the tensile properties.

Acknowledgements The authors wish to thank Mr D. Graham for his assistance in the preparation and heat treatment of the cast plates and Mr M. Ricchetti for assistance with the fatigue testing. This work was conducted under an alliance between the Australian Magnesium Corporation and the Australian Cooperative Research Centre for Cast Metals Manufacturing. The CRC for Cast Metals Manufacturing (CAST) was established under and is supported by the Australian Government’s Cooperative Research Centres Scheme.

References [1] C. J. Bettles, C. T. Forwood, D. H. St.John, M. T. Frost, D. S. Jones, Ma Qian, G-L. Song, J. R. Griffiths and J. F. Nie Magnesium Technology 2003, ed. H.I.Kaplan, 223 (2003). [2] W. Schoeffmann, F. Beste, M. Atzwanger, H. Hick and U. Sauerwein, SAE Technical Paper 2003-01-0406. [3] E. O. Hall, Proc. Phys. Soc. B64, 747 (1951). [4] N. J. Petch, J. Iron Steel Inst. 174, 25 (1953). [5] Ma Qian, D. H. StJohn and M. T. Frost, International Conference on Magnesium Alloys and Their Applications, 2003, Wolfsburg. [6] K. Pettersen and S. Fairchild, SAE Technical Paper 970326. [7] British Standard BS 3518 Methods of Fatigue Testing, Part 5. [8] C. J. Bettles, C. T. Forwood, D. H. St.John, M. T. Frost, D. S. Jones, Ma Qian, G-L. Song, J. R. Griffiths and J. F. Nie, SAE Technical Paper 2003-01-1365. [9] C. J. Bettles, Li Zheng, K. Venkatesan and Ma Qian, International Conference on Magnesium Alloys and Their Applications, 2003, Wolfsburg.

106

Thermal Properties of Mg-Li and Mg-Li-Al Alloys Alexandra Rudajevováa, Pavel Lukáča, Stanislav Kúdelab a b

Department of Metal Physics, Charles University, Prague Institute of Materials and Machine Mechanics SAS, Bratislava

1

Introduction

The magnesium-lithium system is of both theoretical and technical interest. From the equilibrium diagram given in [1], it can be seen that magnesium-lithium alloys exibit two phase structures, the α (hcp) up to 5 wt.% Li and β (bcc) above 11 wt.% Li. Between 5 and 11 wt% Li alloys are composed of two phases α and β. The α phase exhibits a moderate strength and low formability, β phase has better ductility, but a lower creep resistance. The two phase structure (α and β) shows an interesting compromise in propeties because they combine the moderate strength of the α phase and an excellent ductility of the β phase. A superplastic behavior of the two-phase alloys has been reported, too [2]. The Mg-Li alloys are perspective alloys for preparation of the ultra lihgt metal composites. Both thermal and mechanical properties of composites are influenced by the differences of the thermal properties of their components, above all by the thermal expansion and the thermal conductivity. As to authors knowledge the thermal properties of Mg-Li alloys have not yet been studied systematicly. The aim of the present work is to study the thermal diffusivity, thermal conductivity and the thermal expansion of Mg x Li y Al with x = 4, 8, 12 wt.% and y = 0, 3, 5 wt.% as a function of temperature in the temperature o range of 20 to 375 C.

2

Experimental

The linear thermal expansion was measured in argon atmosphere using the Netzsch 402E dilatometer up to 480 °C. The rate of heating was 5 K/min. The thermal expansion samples were 6 mm in diameter and 50 mm long. The accuracy of the apparatus was determined by comparing measured values of the thermal expansion coefficient (CTE) of pure Mg with tabled results. The measurement of the thermal diffusivity was performed in the temperature range from 20 to 300 °C in argon atmosphere using the flash method described elsewhere [3]. The specific heat capacity of all alloys was calculated using the Neumann-Kopp rule. The temperature dependence of the density was calculated from the density measured at 20 °C by weighing the sample in rectified and dried toluene and from respective values of volume thermal expansion data. The thermal conductivity was calculated using the relation λ = acρ

where a is the thermal diffusivity, ρ is the density and c is the specific heat capacity.

(1)

107

3

Results and Discussion

3.1

Binary alloys

The temperature dependence of the thermal diffusivity of the investigated binary alloys is given in Fig. 1. The thermal diffusivity decreases with increasing content of Li in the alloy at all temperatures. The thermal conductivity values determined by equation (1) are presented in Table 1. -1

Table 1. Thermal conductivity of the studied alloys in W m K

-1

25

100

150

200

250

300

Mg4Li Mg8Li Mg12Li

58.7 52.0 48.1

64.6 53.4 49.2

69.1 56.6 50.7

71.5 58.4 52.4

72.0 59.6 52.6

73.5 60.8 54.5

45

Mg4Li Mg8Li Mg12Li

0.4

CTE [10-6 K-1]

thermal diffusivity [10-4 m2 s-1]

Temperature [°C]

0.3

0.2

40

rule mixtures

Mg12Li 35 Mg8Li

30

Mg4Li

25

0

100

200

300

400

temperature [oC]

Figure 1. Temperature dependence of the thermal diffusivity for Mg4Li, Mg8Li and Mg12Li alloys

0

100

200

300

400

temperature [oC]

Figure 2. Temperature dependence of the CTE for Mg4Li, Mg8Li and Mg12Li alloys

The temperature dependence of the CTE for three studied binary alloys is shown in Fig 2. It can be seen that the CTE increases with increasing Li content. It means that the attractive forces or bond energy between atoms decreases when Li content in the alloys increases. While the temperature dependences of the CTE for Mg4Li and Mg12Li are linear, the temperature dependence of the CTE of Mg8Li alloy exhibits a complex character. The concentration dependence of the CTE at room temperature is presented in Fig. 3. Our results are up to 31.4 at.% Li, above this concentration, literature data are shown [4]. It can be seen from this figure that the CTE in single phase ranges changes only a little with concentration. The great changes of the CTE are perceptible in the two phase range. The Mg8Li alloy is a mixture of α and β phases. There is a great difference between the CTEs for both phases (see Fig. 2). If the Mg8Li alloy is cooled from the manufacturing temperature to room temeperature, the individual phases tend to contract unequally, setting up complex internal stresess (tensile character in the high expansion phase and compressive in the low expansion phase) in each crystallite of the alloy. It was found that the CTE of Mg8Li alloy is practically the same for heating and cooling. So it can be assumed that all thermal strains are elastic.

108 thermal diffusivity [10-4 m2 s-1]

50

CTE [10-6K-1]

40

30

20

10

α 0

0.30

0.25

0.20

0

40

100

200

300

400

temperature [oC]

β

α+β 20

Mg4Li3Al Mg4Li5Al

0.35

60

80

100

content of Li [at.%]

Figure 3. Concentration dependence of the CTE, 䊏 our results, 䊐 literature data [4]

3.2

Figure 4. Temperature dependence of the thermal diffusivity for Mg4Li3Al and Mg4Li5Al alloys

Ternary alloys

The temperature dependences of the thermal difussivity of Mg4Li3Al and Mg4Li5Al alloys are shown in Fig.4. It can be seen that the thermal diffusivity increases with increasing temperature up to 225 °C for Mg4Li3Al and 250oC for Mg4Li5Al. Above these temperatures the values of the thermal diffusivity are practically independent of temperature. The phase relations in magnesium-lithium-aluminium alloys were studied by Levinson and McPherson [5], who proposed the isothermal sections of the phase diagram for temperatures 100, 200, 300 and 400 °C. According to these phase diagrams we suppose that both alloys are composed from two phases at room temperature, from the solid solution of Mg-Li and Al and phase AlLi. Above 225 °C and 250 °C the solution of AlLi phase occurs. The solution of this phase is connected with a decrease in the thermal diffusivity. Similar behaviour has been reported for AZ91 alloy [6]. The values of the thermal conductivity for Mg4Li3Al and Mg4Li5Al alloys are given in the Table 2. The values were determined only up to 200 °C because above this temperature the phase transition occurs.

CTE [10-6 K-1]

40

Mg4Li3Al Mg4Li5Al

35

30

25

0

100

200

300

400

temperature [oC]

Figure 5. Temperature dependence of the CTE for Mg4Li3Al and Mg4Li5Al alloys

109 Table 2. Thermal conductivity of Mg4Li3Al and Mg4Li5Al alloys

Temperature [oC]

25

100

150

200

Mg4Li3Al Mg4Li5Al

47.2 42.6

54.1 49.9

58.7 54.5

61.4 57.3

The temperature dependence of the CTE for both ternary alloys is shown in Fig.5. The CTEs increase linearly with increasing temperature up to about 250 °C. Above 250 °C a stronger increase in CTEs with temperature is observed. The increase is connected with the phase transition; with the solution of AlLi phase. This behavior is in agreement with the temperature dependence of the thermal diffusivity.

4

Conclusions

The thermal diffusivity and thermal conductivity of Mg x Li (x = 4, 8, 12) alloy decrease with increasing Li content in the alloys. The thermal conductivity of all alloys increases with increasing temperature. The CTE of the binary alloys studied increases with increasing Li content in the alloys. In the ternary alloys Mg4Li3Al and Mg4Li5Al the solution of AlLi phase occurs at about 250 °C. This phase transition causes a stronger increase in the CTE with temperature and decrease of the thermal diffusivity with temperature.

Acknowledgements The authors acknowledge financial support from the Grant Agency of the Academy of Sciences under Grant A211 2302. S. Kúdela gratefully acknowledge financial support from the Grant Agency of the Slovak Republic VEGA (Project No 2/2078/22).

References [1] Binary Alloy Phase Diagrams, (Ed.: T. B. Massalski), 2nd, ASM International, Materials Park, Ohio, 1996. [2] P. Metenier, G. Gonzales-Doncel, O. A. Ruano, J. Wolfenstine, O. D. Sherby, Mater. Sci. Eng. 1990, A 25, 195. [3] A. Rudajevová, Mat. Res. Bull. 1991, 26, 1363. [4] R. W. Lynch, L. R. Edwards, Appl. Phys. 1970, 41, 5135. [5] D. W. Levinson, D. J. McPherson, Trans. ASM 1956, 46, 689. [6] A. Rudajevová, J. Kiehn,, K. U. Kainer, B. L. Mordike, P. LukáÝ, Scripta Mater. 1999, 40, 57.

110

Designing for Improved Creep Resistance in the RE – Free Magnesium Alloys Vladimir G. Tkachenko International Center for Electronic materials Science and Applied Problems of Aerospace Technology, 3, Krzhyzhanovsky Str, Kyiv, 03680, Ukraine

1

Introduction

Analysis of ongoing magnesium alloy development activities throughout the world indicates current efforts in the development of diecasting alloys with improved elevatedtemperature properties and with acceptable cost to the automobile industry. Nevertheless, the current commercial magnesium alloys of the Mg-Al and Mg-Al-Zn systems making up the bulk of those used in automotive industry do not exhibit creep resistance while maintaining diecastability and room temperature strength. The second group of these alloys of Mg-Al-RE and Mg-Al-Si systems offer borderline improvement in creep resistance but cost and poor diecastability are their main disadvantages. Except for the Mg-AlCa and Mg-Zn-Al-Ca systems other patented experimental alloys also contain the expensive rare-earth additional additives [1, 2]. The sellers of known magnesium alloys containing a series of expensive rare earth elements essentially restricts their application in a mass production for automotive industry, for instance. The less expensive Mg-2…6% Al0.25…5.5% Ca with composition to give Al2Ca precipitation [1] has had diecastability problems including hot-tearing when prototyping an automatic transmission case. A subsequent version of the alloy, Mg-Zn-Al-Ca has had some promise in eliminating diecast problems [3] but the alloy seems to work only in an extremely narrow composition range and also to exhibit variation in properties. It is obvious that more work is needed to develop and modify chemical compositions of these prospective alloying systems. In this connection the development of cheaper, if possible, the RE-free magnesium alloys is required which takes into consideration that in low cost magnesium alloys solid solution – and age – hardening strengthening mechanisms are not enough to decrease time-depended (long-term) creep deformation according to many customers requirements. The main goal of the present research is to develop magnesium alloys of eutectic origin with improved creep resistance and long-term creep fracture strength at higher temperatures (150 °C and above) and higher stresses (64 MPa and above) by the rational allying and related nanoclustering a structure based on knowledge of the thermally activated (rate-controlling) creep mechanisms, to operate in casting (induction melted alloys). New experimental magnesium alloys of Mg-Ba, X and Mg-Al-Ca, X systems to be proposed and their composite materials with dispersions of ceramic components (SiC, AlN, etc.) offer distinct advantages in different dynamical applications including novel branches of automotive, airspace and radioengineering industries where low density and high stiffness (modulus-to-density ratio) as well as low moment of inertia, excellent resistance to timedependent deformation, and dimensional (temperature/temporal) stability are essential.

111

2

Experimental

Discrete temperature spectra of internal friction (inelastic relaxation) in commercially pure casting magnesium alloys were measured by advanced mechanical spectroscopy technique using semi-automatic relaxometer at audio-frequency (~1 Hz) [4] in order to perform thermo-activated analysis and to evaluate the mechanisms of resonance absorption of elastic energy by structural defects at elevated temperatures. According to Conrad concept [5] activation parameters for creep deformation needed to identify its ratecontrolling mechanism can be derivable from macroscopic mechanical property data, i.e. macrodeformation partials such as ∂ ln ε& / ∂(1/ T ) and ∂ ln ε& / ∂σ if activation energy (a.e.) as function of stress independent of temperature. The experimental technique to be used for constant structure creep tests has been described in details elsewhere [4].

3

Results

Development of higher creep-resistant alloy is a complicated task and requires an understanding of the basic principles of creep resistance as well as the mechanisms of creep in magnesium alloys. Some good work is underway along these lines to study the low temperature strengthening mechanisms of commercial alloys (AZ91D, AE42, AS41) but the science of magnesium creep resistance in diecasting alloys is still in its infancy [3]. Moreover, future activity is need to be based on an industry-wide effort to create this science as alloy development activity goes further. We have recently taken up this challenge by targeting the development of a new advanced magnesium alloys with cluster-forming structure [4] based on the Mg-Al-Ca ternary systems with superior compression and tensile creep resistance at 150 °C under stress 64 MPa which are currently undergoing patent application process and have acceptable cost to the automotive industry. The recent discovery of chemically combined and mobile pair defects, i.e. primary nanoclusters consisting of solute alloying element (AE) and excess vacancy (EV) with certain bonding energy in as-cast magnesium alloys of Mg-Ba system and as-cast aluminium alloys of Al-Li system [5–7] is confirmed first of all by mechanical spectroscopy (Fig. 1) and also results from analyzing the X-ray photoelectron and Auger electron spectra [6, 7]. Figure 1. The typical discrete temperature In the Auger-electron spectrum of the Mg-Ba alloy spectra of internal friction Q (T) for there have observed chemical (binding energy) schift casting magnesium alloy of Mg-Ba, Al (about of 2 eV/atom) due to alloying of the Mgnsystem (nanocluster assembly mechacluster with barium to form chemical bonding in this nism). ‘immiscible’ system (Fig. 2). At the most early stages microalloying by selected elements causes diffusive nanoclustering of solid solution to be extended near the grain boundaries (GBs). The process is then accompanied by activation of nanocluster assembly GB strengthening mechanism in α-matrix of eutectics. -1

112 The effect of inducing the excess (extrinsic, structural) EV in the field of mobile solutes stresses was predicted theoretically and then explained in a series of our works [4 – 7]. It was shown that increasing of the solute-induced concentration of EV changes vacancytype mechanism of diffusion and consequently, the concentration dependence of self-diffusivity. In accordance with this, observed diffusion by mixed (nanocluster) mechanism is considered to be one of the most probable rate-controlling nanocluster mechanism which Figure 2. Chemical shift in low – (i) causes the anomalous deviation of solute diffusivity energy region of the Auger – electron from classical (Arrhenius) behavior and from Vegard spectrum (LVV – transition) for law for solid solution, and (ii) is responsible for creep Magnesium (a) and Mg-1% Ba alloy resistance al elevated and higher temperatures. (b) while alloying the embedded It would be reasonably to suppose that creep resisclusters Mg with barium Ba . tance and creep strength of magnesium alloys based on Mg-Al-Ca, X system can be further improved by alloying additions which form favourable cluster-forming structure to be characterized by distribution of AE bound with EV and gives rise to delayed creep yield/fracture due to nanoclusters assembly near the GB denuded zones and to strengthening phase transformation at the GB. Thus, elevated-temperature GB strengthening is thermally activated process which is controlled by diffusive nanoclustering a structure near the GB and thermal segregation of AE to the GB. This statement is confirmed by measuring the damping components of inelastic relaxation in substitutional solid solution of Mg-Ba, Al system (Fig. A). The discrete temperature spectrum of internal friction Q-1 (T) reveals relaxation maximum (A) of vacancy origin with a.e. of about 0.8 eV at 310…345K to be caused by pair relaxation, i.e. EV-AE nanoclusters reorientation (curve 1 – quenching) and GB relaxation peak (B) to be characterized by GB sliding. Thermally activated recovery of the clusterforming structure is accompanied by suppressing the GB-induced relaxation (B) with a.e of 1.1 –1.3 eV at 450 –475 K and appearing (peak C) the segregation of AE at the GB with a.e. of about 1.5eV at 520 –565 K (curve 2 – annealing, 3 – slow cooling). The process is completed by blocking GB dislocation sliding due to both strengthening Figure 3. Summarized creep data for the as-cast Mg mechanism of nanocluster assembly near the AZ91D alloy and magnesium matrix composite GB and phase transformation at the GB. material of Mg-Al-Ca-SiC casting system at conThe effect of their reinforcing influence after stant creep temperature and constant creep strength. nanocluster assembly into microclusters by chemical reordering and relaxation mechanism have stimulated considerabe interest in the fundamental nature of the process [7] and will open a promising field for the synthesis of new classes of metallic light alloying systems of eutectic origin as well as immucsible systems such as Mg-Ti, Mg-Cr etc. Creep strength, freedom from cracking can be further n

m

113 improved by dispersing and strengthening the phase components as well as enhancing the contribution of eutectic strengthening (Fig. 3). The activation energy for creep deformation in proposed magnesium alloys of Mg-Al, Ca system is found to be equal to that for the GBs nanocluster solute diffusion-accommodated the GBs sliding (Fig. 4a). For the rate-controlling creep mechanism steady state creep rate ( ε&GB ) proportional to an effective diffusion coefficient (Deff) for slower – diffusing AE can be described by the following relation:

Figure 4. (a) Stress dependence of the steady state strain rate (with stress exponent n) and (b) Arrhenius plot in the power law creep regime for commercial cast Mg AZ941D (1) and cast alloy of Mg-Al-Ca, X system with cluster-forming structure (2).

D ⎛ σ ⎞ ε&GB ≅ K GB ⋅ reff2 ⋅ f ⎜ ⎟ b ⎝ EGB ⎠

r where KGB is constant, b –Burgers vector; ⎛ σ ⎞ f⎜ ⎟ is a yet unspecified function of the GBs ⎝ EGB ⎠ modulus-compensated stress [8]. Introducing the additions of Ca increases AE of the GBs diffusion from 0.2– 0.4 eV up to 0.6 –0.8 eV. Straight line in the theoretical coordinates ln ε& vs ln σ (Fig. 4) suggests the power relationship of the form ε& ≅ σ n , where stress exponent n decreases from n = 3.0 (for AZ91D Mg alloy) to n = 1.1 (for new Mg-Al, Ca, X alloy). The results obtained for alloys with cluster-forming structure predict unusual the GBs creep behavior by the Coble-type diffusion mechanism. This means that clusters as thermal barriers promote to reducing the creep rate in a wide temperature range i.e. the GBs diffusion mobility of AE is lowed in the alloying systems with cluster-forming structure. Therefore creep rates of new experimental alloys with cluster-forming structure are essentially slower than that for commercially available AZ91D (Dow Chemical Corp., USA) (Fig. 4).

4

Discussion

4.1

Dynamical clustering structure concept development

Beyond all doubt, microstructure of as-cast pure magnesium alloys of eutectic origin is supersaturated by excess vacancies [5] generated by stress fields of mobile impurities such as hydrogen and the like [6]. The vacancies are trapped by alloying elements so that the presence of a heterogenous nucleation such as the GBs is necessary for the pair defect reaction to proceed. Therefore the thermally activated nanoclusters of EV-AE formed by rapid (short-range) mechanism [5–7] segregate as faster-diffusing components towards the GBs producing (by nanoclusters assembly mechanism) alloyed microclusters of Mgn (AE)m in denuded, free of participates near-the GBs zones where recombination of EV is possible. Nanoclusters can relax sufficiently to reach a stable microcluster structure with

114 low energy. Thus, microclusters should be regarded as thermally stable separations with intermetallide-like ordering or pre-nanophases with saturated (non-stoichiometric) composition near-by to corresponding chemical compound. In particular, in Mg-Ba alloys stable MgnBam cluster configuration consisting of 16 atoms of Mg is limited by separation of Mg16Ba2 (against stoichiometric intermetallide of Mg17Ba2 [4]). Based on dislocation damping data in a wide temperature range [4] we state that at higher temperatures and higher stresses the creep mechanism in new experimental magnesium alloys is controlled by excess vacancy – alloying element nanoclusters diffusion flux (to the GBs) rather than thermally-induced dislocation migration. 4.2

Electronic structure of the cluster

The results obtained are consistent with the fact that microclusters responsible for intermatallide-like ordering could be treated as nuclei of the intermatallide phase (Mg17Ba2 in Mg-Ba system, MgAl12-XCaX in Mg-Al-Ca system, etc.) which are characterized by concomitant evolution of the cluster-induced electronic structure inside solid solution extended nearby the GBs. This is confirmed by using the Auger-electron – and X-ray photoelectron spectra measurement techniques have proved to be very powerful tool in elucidating the cluster-induced electronic structure in solid solutions and eutectic alloys of Al-Li and Mg-Ba systems [7]. The coincidence of the experimental X-ray photoelectron spectrum for Mg-Ba alloys and the theoretical spectrum calculated from the first principles for MgnBam – embedded clusters [7] indicates existence of cluster size – dependent and strong electronic structure, the binding energy shift of which can be interpreted as a cohesive energy gain (per atom) increasing with cluster size (from the monomer, monosized cluster to a bulk microcluster) and in this way enhancing the elastic (Young’s modulus) and creep strength properties of the above systems near-by the GBs. The bonding d-state of Barium is dominant around the Fermi level hybridizated with s-, and p-states of magnesium. 4.3

The GBs strengthening (rate-controlling) creep mechanism

Thus cluster-dependent creep strength is closely related to the concomitant evolution of the cluster-induced electronic structure which should take into account by a complete theory of solute diffusion and dislocation theory in terms of diffusion/dislocation relative contributions. First, the resistance to the lattice dislocations neigbouring the GBs increases during their motion through the thermally stable and ordered microclusters with hybride electronic structure more strong than that of α-Mg matrix. Second, the ability of size-seleted cluster ordering to relax energy and suppess a premature embrittlement is also related to the kinetics (thermally – induced migration) of the slower-diffusing component which facilitates developing the delayed (thermally-activated, subcritical) fracture at the vicinity of GBs and is an important factor in determing the rate-controlling creep mechanism [4]. After vacancy recombination agglomerated clusters form more stable configurations which act as closed domains keeping their unique indivisible properties at higher temperatures. The participates play an important role on pinning the GBs and reveal the direct influence on stabilizing their structure. For example, addition of Ca into Mg-Al alloy can result in precipitation hardening of the GBs that keeps the pinning effect up to higher temperatures due to the ternary MgAl12-XCaX formation. Besides, introducing the non-metallic particles of SiC in the nanoregime additionally strengthens the GBs of α-Mg even in a small amounts.

115 Thus, the strengthening of near-by the GBs internal interfaces by cluster mechanism in combination with the GBs strengthening by precipitate/dispersion hardening mechanism results in improving the creep resistance and the long-term creep strength of new experimental magnesium alloys at higher creep and larger creep stresses (Fig. 3 and Fig. 4). In order to account for these findings and the double strengthening effect a structural interface model has been developed [7] which is consistent the fact that in high – creep resistant magnesium alloys quasi-brittle and sudden failur is overcome by forming the initial crack – absorbing matrices in eutectic compositions and particulate-reinforced magnesium matrix composites by means of activation of inhibition mechanisms which allow delayed fracture in a wide range of creep stresses.

Concluding Remarks 1. There are at least three ways to decrease the GBs diffusion creep rate of magnesium at higher temperatures and higher stresses. They are: (i) alloying with segregated additions increasing the Young’s elastic modulus of the GBs; (ii) clustering a crystal structure near the GBs and its chemical bond strengthening more resistant to thermally-induced diffusion and moving dislocation; (iii) grain boundary structural and phase transformations induced by solute segregations to block the mobility and multiplication of the GBs dislocations. 2. Alloying of the immiscrible systems (Mg-Ba, etc) is found to be occur by alloying microclusters Mgn embedded inside solid solution extended near-by the GBs/interfaces. 3. In eutectic alloys and immiscrible systems cluster-forming structure is responsible for rate of chemical reactions and the nanocluster-microcluster transience of separations.

References [1] Pekguleryuz, M. O. and Luo, A, International Patent Application WO 96/25529, 1996. [2]. Luo, A. and Sinoda, T, US Patent US 5855697, 1997. [3] Proc. of Conf. on Magnesium Alloys and their Application (Eds.: B. L. Mordike and K. U. Kainer) in the Verstoff Information geselleschaft, Wolsburg, Germany, 1998, 37–47. [4] V. G. Tkachenko, Strength Physics of Less Common Metals and their Alloys, 2nd. ed., Cambridge Sci. Publish., London, 2002 (ISBN-189 832 6576); Nauk. D., Kiev, Ukraine, 1996, p. 359 (in Rus.). [5] V. G. Tkachenko, B. G. Strogin, I. N. Maksimchuk, V. V. Friezel, V. A. Grecko, S. P. Likhtarovich, Int. J. Hydrogen Energy. 1996, 21, 1091–1096. [6] V. G. Tkachenko, V. A. Tatarenko, I. I. Schuljak, A. N. Malka, A. M. Strutinsky, Metallofiz. Noveishei Teknol., 2001, 23, 361–386. [7] V. G. Tkachenko, I. I. Schuljak, A. M. Strutinsky, V. V. Friezel, O. V. Crivik, in the Hydrogen Materials Science and Chemistry Metal Hybrides (Eds.: N. Vezigrogly et al.), Kluver Acad. Publish., the Netherlands, 2002, 77–95. [8] O. D. Sherby, R. H. Klundt, and A. K. Miller, Met. Trans., 8A, 843–850.

116

Ageing Characteristics and Creep Resistance of Mg-Y-Nd-Sc-Mn Alloy Ivana Stulíková,1), 2) Bohumil Smola,1), 2) Jitka Pelcová,1) Barry L. Mordike2), 3) 1)

Faculty of Mathematics and Physics, Charles University, Prague, Zentrum für Funktionswerkstoffe gGmbH, Clausthal-Zellerfeld, 3) Institute of Material Science and Engineering, Technical University, Clausthal-Zellerfeld 2)

1

Introduction

High thermal stability of precipitated phases is decisive for a reasonable stability of mechanical properties and for a good creep resistance up to high temperatures. Not only a volume fraction but also arrangement, orientation relationship and aspect ratio of precipitates affect mechanical and creep properties [1, 2]. It was proved that prismatic precipitate plates in a triangular arrangement play an important role in strengthening of WE54 alloy by providing effective barriers to gliding dislocations [3]. Both β1 transition phase and the equilibrium phase β form as plates on { 10 10 } plane of the α-Mg matrix [3, 4, 5]. The peak hardening of more than 50% of the initial value in Mg-15 wt.%Gd alloy is achieved by the precipitation of a metastable β´ phase in the form of fine plates parallel to all three { 1120 } α-Mg matrix planes. The equilibrium Mg5Gd phase precipitates form a triangular arrangement of plates with the same orientation relationship to the α-Mg matrix as the WE54 β phase [6]. The efficiency of this precipitate structure diminishes at temperatures above 250 °C due to its coarsening and/or the increasing role of cross slip of basal dislocations and non-basal slip. It was recognized that Mn2Sc phase precipitating in fine discs parallel to basal planes ensures a high creep resistance in Mg-Sc-Mn, Mg-Y-Sc-Mn or Mg-Gd-Sc-Mn alloys [7, 8, 9]. The dense dispersion of basal discs of Mn2Sc incoherent phase (diameter ≈ 1015 nm, thickness ≈ 3 nm) in grain interiors persists during a long exposure even at 350 °C without any coarsening. Nevertheless, to obtain a significant age hardening the relatively high content of rare earth is necessary (≥ 10 wt.%Gd). To utilize the lowering effect of Nd addition on Y solubility in Mg known from the WE-type alloys, which should improve age hardening, and to maintain the creep resistance at high temperatures the Mg-Y-Nd-Sc-Mn alloy was squeeze cast. The aim of this contribution is to report phase transformations due to thermal treatment and the effect of phases precipitated on creep behavior and mechanical properties.

2

Experimental

The alloy MgY4Nd2Sc1Mn1 containing 3.71 wt.% Y, 2.12 wt.% Nd, 1.28 wt.% Sc and 1.18 wt.% Mn was squeeze cast under a protective gas atmosphere (Ar + 1%SF6). The microstructure is typical for a cast structure with cellular grains ≈ 70 μm, which are separated by a grain boundary eutectic consisting of α-Mg matrix and β phase as observed in

117 WE54. The alloy has very little porosity. Its density was determined to be 1844 kg/m3 as compared to the calculated 1837 kg/m3. The isothermal and isochronal annealing response of relative electrical resistivity changes was determined in the range 20 °C to 530 °C. Isothermal as well as isochronal annealing carried out in steps of 30 °C/30 min were followed by quenching. Heat treatment was performed in a stirred oil bath up to 240 oC or in a furnace with argon protective atmosphere at higher temperatures. The H-shaped specimens machined to dimensions of 2 x 6 x 80 mm3 were used for resistivity measurements at 77 K after each heating step. The value of the length represents the gauge length for the resistivity measurements. Relative electrical resistivity changes Δρ/ρ were obtained within an accuracy of 10-4. The resistivity was measured by means of the dc four-point method with a dummy specimen in series. The influence of parasitic thermoelectromotive force was suppressed by current reversal. The thermal stability of mechanical properties during described heat treatment was measured by Vickers hardness HV3 at room temperature. The tensile tests were performed at temperatures from room temperature to 400 °C at a constant strain rate ≈ 10-4 s-1. The minimum creep rates were determined (mostly by creep dip tests) at 225 °C - 350 °C at stress 30 or 40 MPa. Transmission electron microscopy (TEM) and electron diffraction were carried out to determine the microstructure of the alloy after specified treatments. An analysis of the phases precipitated out was also supported by energy-dispersive-X ray microanalysis. Studies of the microstructure were carried out on JEOL JEM 2000FX and Philips 200CM electron microscopes equipped with LINK AN 1000 and EDAX microanalyser, respectively. The specimens for TEM were prepared by the same isochronal or isothermal annealing procedure as those for electrical resistivity and hardness measurements.

3

Results and Discussion 80

HV3

Fig. 1 shows the response of HV3 to the isothermal annealing at 500°C for as 75 cast material. It does not prove the possibility of T4 treatment. TEM studies 70 showed that the GB eutectic dissolves during this 65 annealing, but the precipitation of Mn2Sc phase occurs already after 1 hour 60 of annealing. The particles 0 2 4 6 8 10 of Mn2Sc phase in the time [h] form of fine discs precipiFigure 1. Annealing response of HV3 to isothermal annealing at 500 °C. tated parallel to basal planes of the matrix and they coarsen continuously (≈ 60 nm after 1 hour, ≈ 100 nm after 8 hours of annealing). The impossibility of a homogenisation or only very narrow temperature and composition ranges, where a homogenisation is possible, were reported also in other Sc and Mn containing Mg-R.E. alloys [7, 10, 11]. MgY4Nd2Sc1Mn1 - 500°C

118 20

Δρ/ρ0 [%]

The electrical resistivity response of as cast alloy 15 on the isochronal step-byMgY4Nd2Sc1Mn1 as cast step annealing up to 10 570 °C shows a slight 5 decrease followed by an increase of Δρ/ρ0 exceed0 ing the initial value (90 °C - 240 °C) – Fig. 2. The -5 main feature of this curve is the decrease of electri-10 0 100 200 300 400 500 T[°C] 600 cal resistivity between 240 °C and 360 °C. The Figure 2. Isochronal annealing curve of relative electrical resistivity changes main precipitation procΔρ/ρ measured at 77 K. ρ is the value in as cast state. esses can be expected in this temperature range. The electrical resistivity increase in the temperature range 360 °C – 570 °C indicates dissolution of some precipitated phases and of GB eutectic into the matrix. TEM revealed the precipitation of the c-based centered orthorhombic (cbco) transition phase in the form of fine plates near the GB eutectic after the annealing up to 270 °C – Fig.3. Microstructure and orientation of cbco phase to the matrix (plates parallel to all three { 1120 } α-Mg matrix planes) are the same as those of the metastable β´ phase in Mg-Gd alloy [6]. This phase transforms into a stable phase of Mg5Gd-type (fcc structure, a = 2.23 nm) as observed after the annealing up to 360 °C. The plates of the stable phase 0

0

Figure 3. ED pattern of the alloy annealed up to 270 °C. Pole [0001]Mg with all three possible [001]cbco patterns.

Figure 4. Mn2Sc basal discs, thin hexagonal plates parallel to (0001)Mg plane and stable β phase plates parallel to ( 1010 )Mg plane in the alloy annealed up to 360 °C.

119

HV3

(diameter ≈ 500 nm) are parallel to { 10 10 } matrix planes. Mn2Sc incoherent phase in the form of 90 fine discs (diameter ≈ 20 nm, thickness ≈ 5nm) parallel to matrix basal planes was observed after 80 annealing up to 360 °C. Moreover, very thin basal plates of a metasta70 as cast -200°C ble hexagonal phase containing hom. 500°C/8 h. - 200°C as cast-250°C rare earths and Mn were observed, 60 too – Fig. 4. The same phase was 0 10 20 30 40 50 60 t [h] observed in other Mg-rare earth Figure 5. HV3 hardness response of MgY4Nd2Sc1Mn1 to iso- alloys with Sc and Mn [9]. No phases containing Y or Nd were thermal annealing at 200 °C and 250 °C confirmed after step-by-step annealing up to 480 °C, the Mn2Sc phase particles coarsened to the size of 40-100 nm. The yield stress and tensile strength of the alloy in as cast condition are comparable to or better than those of the as cast WE54 alloy [12] in the temperature range 20 °C-400 °C. The hardness HV3 showed the best response after the annealing at 200 °C – Fig. 5. The most appropriate T5 heat treatment was determined to be 200 °C/40 h. If the alloy is preheated at 500 °C/8 h, the age hardening at 200 °C is more pronounced due to the higher over-saturation of the matrix with Y and Nd from the GB eutectic dissolving during the high temperature treatment. Transition cbco phase precipitates during the T5 annealing in dense triangular arrangement of prismatic plates (diameter ≈ 40 nm, thickness ≈ 5 nm) along GB eutectic (the thickness ≈ 3 μm). The plates precipitate only on dislocations inside the grains 100

MgY4Nd2Sc1Mn1

Figure 6. Structure of the alloy after T5 annealing. Layer of dense cbco phase along GB eutectic.

Figure 7. Structure of the alloy after annealing at 250 °C for 45 h. Note the similarity to Figure 4.

120 350

325

300

275

250

225

200

T [°C]

minimum creep rate [s-1]

– Fig. 6. The microstructure of as cast alloy annealed 45 hours at 1.E-06 MgY4Nd2Sc1Mn1 as cast σ = 40 MPa 250 °C is the same as after the MgY4Nd2Mn1Sc1 T5 1.E-07 isochronal step-by-step annealing of as cast state up to 360 °C. Both 1.E-08 the plates of the stable phase iso1.E-09 morphous with Mg5Gd and the discs of the stable Mn2Sc incoher1.E-10 ent phase were identified as well 1.E-11 as the fine basal plates of a me0.0016 0.0017 0.0018 0.0019 0.002 0.0021 0.0022 reciprocal temperature [K ] tastable R.E.- and Mn- containing Figure 8. Temperature dependence of minimum creep rates at 40 phase – Fig. 7. MPa Fig. 8 shows the temperature dependence of minimum creep rate of as cast and T5 treated alloy compared to as cast Mg - 6wt. %Y - 3wt. %Nd - 0.6wt. %Zr (WE63) alloy crept at 40 MPa. The alloy is superior to the WE-type alloy by at least one order of magnitude of minimum creep rates. The lowering of minimum creep rate by increasing temperature from 250 °C to 275 °C is most probably caused by simultaneous precipitation of Mn2Sc basal discs during the creep test. One can not exclude also the precipitation of R.E.- and Mn- containing metastable phase. 1.E-05

WE63 as cast

-1

elongation [%]

2.5 WE43 as cast

350°C, 30 MPa

2

1.5

1

MgY4Nd2Sc1Mn1 - as cast

MgY4Nd2Sc1Mn1 - T5

0.5

0 0.E+00 5.E+04 1.E+05 2.E+05 2.E+05 3.E+05 3.E+05 4.E+05 4.E+05 5.E+05 5.E+05 time [s]

The excellent creep resistance of this alloy, particularly in the T5 condition, persists up to high temperatures – see creep curves at 350 °C and 30 MPa in Fig. 9 compared to WE43 alloy. It demonstrates the efficiency of the Mn2Sc phase basal discs in the dense distribution together with the triangular arrangement of stable β phase plates as obstacles for creep deformation.

Figure 9. Creep curves at 350 °C and 30 MPa

4

Conclusions

The possibility of T4 or T6 heat treatments was not proved for MgY4Nd2Sc1Mn1 alloy. Although the age hardening at 200 °C after a preheating of the alloy at 500 °C is much more pronounced, the Mn2Sc phase coarsens during the annealing at 500 °C. The transition cbco phase known from Mg-Gd system as well as from WE alloys precipitated in a dense triangular arrangement of prismatic plates near the GB eutectic during T5 treatment. The plates precipitate only on dislocations inside the grains. After annealing at higher temperatures (250 °C for a long time or up to 360 °C isochronally) both the plates of the stable phase isomorphous with Mg5Gd and the discs of the stable Mn2Sc incoherent phase were identified as well as the fine basal plates of a metastable R.E.- and Mn-containing

121 phase. Even though the metastable phases disappear at high temperatures, the studied alloy, especially in the T5 condition, is superior to WE43 alloy in the creep resistance at least up to 350 °C.

Acknowledgement The support by the German Research Foundation (DFG), by the GACR (# 106/03/0903) and by the Ministry of Education of the Czech Republic within the framework of the research program MSM 113200002 are gratefully acknowledged.

References [1] J. F. Nie, B.C. Muddle, in Mg Alloys and Their Applications (Eds: B.L. Mordike, K.U. Kainer), Werkstoff-Informationsgesellschaft, Frankfurt, 1998, p. 169. [2] J. F. Nie, Scr. Mater. 2003, 48, 1009-1015. [3] J. F. Nie, B.C. Muddle, Acta mater. 2000, 48, 1691-1703. [4] G.W. Lorimer, in Proc. London Conf. Magnesium Technology (Ed. H. Baker), Inst. of Metals, London, 1986, p. 47. [5] I. Stulikova, P. Vostry, B.L. Mordike, B. Smola, in Progress on Heat Treatment of Materials, IFHT, Budapest, 1986, p. 1993. [6] P. Vostry, B. Smola, I. Stulikova, F. von Buch, B.L. Mordike, phys. stat. sol (a) 1999, 175, 491-500. [7] B. Smola , I. Stulíková, J. Pelcová, F. von Buch, B.L. Mordike, phys. stat. sol. (a) 2002, 191, 305-316. [8] I. Stulikova, B. Smola, F. von Buch, B.L. Mordike, Mat.-wiss. u. Werkstofftech. 2003, 34, 102-108. [9] B. Smola, I. Stulikova, J. Pelcova, F. von Buch, B.L. Mordike, Z. Metallkd. 2003, 94, 553-558. [10] R. Schmidt-Fetzer, J. Gröbner, D. Kevorkov, Mat.-wiss. u. Werkstofftech. 2001, 32,13-19. [11] J. Gröbner, D. Kevorkov, R. Schmidt-Fetzer, in: Magnesium Technology 2001 (Ed. J. Hryn), TMS, Warrendale, 2001, p. 105. [12] I. Stulikova, B. Smola, B.L. Mordike, phys. stat. sol. (a), 2002, 190, R5-R7.

122

Deformation Behaviour of Mg-Li-Al Alloys at Room and Elevated Temperatures Zdeněk Drozd, Zuzanka Trojanová, Viera Gärtnerová Faculty of Mathematics and Physics, Charles University, Prague

Introduction Magnesium-lithium alloys are the lightest metallic materials that can be used for structural purposes. Alloying by lithium decreases the density of the alloy from 1.74 g.cm-3 (pure magnesium) to 1.35 g ⋅ cm-3. The solubility of Li in Mg is about 5.5 wt% (17 at%). The alloys with the content of Li higher than 11.5 wt% crystallize in bcc phase (β phase) and the solid solution of Mg with low Li content up to 5.5 wt% (α phase) has hcp structure (as Mg pure). When the content of Li is between 5.5 and 11.5 wt% there is a mixture of hcp and bcc (α + β) phases in the material. Mechanical properties of the alloys with hcp structure are worse then the ones with bcc structure. Bcc alloys are better for machining and welding then hcp materials, but they have low corrosion resistance, that decreases with increasing content of Li. Therefore the material with the Li content between 5.5 and 11.5 wt% (α + β) could be a compromise with acceptable mechanical properties as well as corrosion resistance. To improve the mechanical properties ternary alloys have been developed. Adding of the elements as Cd, Al, Zn or Si lead to precipitation of coherent transitive phase θ with composition of MgLi2X, where X is the third element. When Al is used as the third element then the α-MgLi phase can be hardened by the precipitation of stable AlLi and Mg17Al12 phases but such hardening is less effective than that achieved in β-MgLiAl where Li prohibits micro-segregation of Al and precipitation of MgAl eutectics [1]. The aim of present paper is to investigate mechanical properties of three Mg-Li-Al alloys (Mg4Li3Al, Mg8Li3Al and Mg12Li3Al) at room temperatures and elevated temperatures up to 200 °C.

Experimental Procedure Cylindrical specimens each of 5 mm diameter and 7 mm lenght were deformed in compression using an Instron machine. Experiments were performed at temperatures between room temperature and 200 °C. The initial strain rate was 2.4 x 10-4 s-1. Stress relaxation tests were carried out at temperature 14 °C that is a little bit below the room temperature. This temperature was achieved using water cooling. The argon atmosphere was used as a protecting atmosphere in the furnace at elevated temperatures.

123

Experimental Results and Discussion Two examples of the microstructure of the studied alloys are introduced in Figures 1 and 2. Both photographs were obtained by optical microscopy. Specimens were etched using Glycol for 30 s and 1% Nital for 15 s. Figure 1 shows the microstructure of the Mg4Li3Al alloy. Eutectics can be well seen in the photograph. Figure 2 shows the microstructure of the Mg8Li3Al alloy. Grain boundary and eutectics are visible there.

Figure 1: Microstructure of Mg4Li3Al

Figure 2: Microstructure of Mg8Li3Al

Stress-strain curves for Mg8Li3Al alloy deformed at various temperatures are introduced in Figure 3. It can be seen from the figure that the deformation stresses at elevated temperatures are significantly lower than the ones at the room temperature. The stress rapidly decreases with increasing temperature. This decrease can be better seen from Figure 4 where the temperature dependencies of the yield stress σ02 and the maximum stress σmax are shown. The stress-strain curves for Mg4Li3Al and Mg12Li3Al were obtain by the same way. The differences between the maximum stresses and the yield stresses for all three alloys at various temperatures can be seen from Figure 5. The differences of those stresses exhibit 300

RT

200 70°C

100

σ02 σmax

σ [ MPa ]

σ [ MPa ]

300

200 100

100°C 200°C

0 0.0

0.1

0.2

0.3

0.4

ε

Figure 3: True stress-true strain curves for Mg8Li3Al alloy obtained in compression tests.

0 0

100

200

T [ °C ]

Figure 4: The dependencies of σ02 and σmax on the temperature for Mg8Li3Al.

124 rapid decrease with increasing temperature. Observed decrease of 200 the difference between the yield stress and the maximum stress with the temperature is more significant for alloys containing bcc structure. 100 Thermally activated slip of dislocations is more easy at higher temperatures. Recovery processes in hcp Mg alloys AZ91, AE42 and AS21 0 were studied by K. Mathis et al [2]. 0 100 200 They found that the dynamical recovery process, including conT [ °C ] servative slip (cross slip) may Figure 5: Differences between the maximum and yield stress compensate the hardening process. Then the maximum flow stress in the hcp alloy Mg4Li3Al should decrease with temperature as the stress necessary for the cross slip in the form

Δσ [ MPa ]

Mg4Li3Al Mg8Li3Al Mg12Li3Al

⎛σ ln ⎜ max ⎝ σ 0,max

⎞ ⎟ = − K (γ sf , ε& ) , ⎠

(1)

where σ0,max is a constant and K (γ sf , ε& ) is function of the stacking fault energy γs,f and the strain rate ε& . In order to estimate parameters of a possible thermally activated process the stress relaxation technique was used. The individual relaxation tests took 300 s. Theoretical relationship between stress decrease rate ( −σ& ) and the stress σ can be expressed in the following form [3]: ln ( −σ& ) = C + n ln σ ,

(2)

where C is a constant and n is the parameter of stress sensitivity ⎛ d ln ε& ⎞ n=⎜ ⎟ . ⎝ d ln σ ⎠T

(3)

Parameter of the stress sensitivity n may be experimentally determined as the slope of the dependence ln ( −σ& ) vs. ln(σ/σ0), where σ0 is the starting stress of the stress relaxation test. The experimental activation volume Va depends on the parameter n by the following expression nkT = σ Va ,

(4)

where k is the Boltzmann constant and T is the absolute temperature. Dependence (2), i.e. dependence of the velocity of the stress drops on the reduced stress for Mg8Li3Al and Mg12Li3Al at the temperature 14 °C, is plotted in Figure 6 and Figure 7, where σ0 is the initial (starting) stress of the relaxation. Parameters n at the beginning and at the end of the relaxation test were determined from these plots. The activation volume was then determined from the relation (4).

125 101

101

σ0=167.9 MPa

σ0 = 161.2 MPa

σ0=194.9 MPa

σ0 = 224.3 MPa

-1

100

-σ [ MPa.s ]

-1

-σ [ MPa.s ]

σ0 = 188.6 MPa σ0 = 244.5 MPa σ0 = 265.9 MPa

10-1

σ0=227.2 MPa

100

σ0=258.9 MPa σ0=282.6 MPa

10-1

10-2

10-2 0.7

0.8

0.9

0.6

1.0

0.7

0.8

0.9

1.0

σ / σ0

σ / σ0

Figure 6: SR curves obtained at various starting stresses at 14 °C (Mg8Li3Al).

Figure 7: SR curves obtained for various starting stresses at 14 °C (Mg12Li3Al)

The values of the stress sensitivity parameters n1 (at the beginning of the SR test) and n2 (at the end of it) are introduced in Table 1 and Table 2 as well as the calculated values of the activation volumes V1, V2 (also for the begin and the end of each SR test). The values σ1 and σ2 in the tables mean the initial stress or the finite stress, respectively. From Figure 8 it can be seen that the activation volume decreases with increasing starting stress of the SR test. The decrease of the activation volume with the flow stress is -28 typical for metallic polycrystals. The values of the activation volume in the order of 10 are typical too. The apparent (experimentally estimated) activation volume Va is proportional to the dislocation activation volume Vd by the relationship Vd = ψ Va ,

(5)

where ψ is the Taylor factor. The values of the dislocation activation volume at the beginning of the SR tests for Mg12Li3Al are introduced in the Table 2. Geometrical interpretation of the activation volume is done by the expression Vd = bdl ,

(6)

where b is the Burgers vector, b is the effective width of an obstacle and l is the length of the dislocation segment pinned at obstacles. Table 1. Results of the stress relaxation (SR) tests obtained at 14 °C for Mg8Li3Al alloy

SR1 SR2 SR3 SR4 SR5

σ1 [ MPa ] 161.2 188.6 224.3 244.5 265.9

σ2 [ MPa ] 128.7 148.5 174.7 188.9 205.2

n1 40.5 35.4 31.3 30.6 28.5

n2

V1 [10-28 m3]

V2 [10-28 m3]

13.6 12.0 11.0 9.8 11.1

10.0 7.4 5.5 5.0 4.3

4.2 3.2 2.5 2.1 2.1

126 Table 2. Results of the stress relaxation (SR) tests obtained at 14 °C for Mg12Li3Al alloy

SR1 SR2 SR3 SR4 SR5

σ1 [ MPa ] 167.9 194.9 227.2 258.9 282.6

σ2 [ MPa ] 114.4 128.8 148.2 168.2 181.7

n1 21.1 18.1 16.3 16.1 15.4

n2 7.3 7.2 7.4 7.3 6.4

V1 [10-28 m3] 4.5 3.3 2.5 2.2 1.9

V2 [10-28 m3] 2.3 2.0 1.8 1.5 1.3

Vd [ b 3 ] 31.5 23.1 17.5 15.4 11.2

12 Mg8Li3Al Mg12Li3Al

[m ]

3

6

V x 10

8

-28

10

4 2 0 150

200

250

300

σ0 [ MPa ]

Figure 8: Stress dependencies of the activation volume for temperature 14 °C.

The velocity of a screw dislocation in a bcc structure is given approximately by [4] v = v0

b L ⎛ ΔH ⎞ ⋅ ⋅ x exp ⎜ − ⎟, lc lc ⎝ kT ⎠

(7)

where v0 is the Debye frequency, v0 b l c is the frequency of vibration of the dislocation with wavelength l c , L l c is the number of activation sides in a free length of dislocation, and x is the distance traveled by the screw dislocation in the glissile form becoming again sessile. ΔH is the activation enthalpy of the glissile – to – sessile transformation and constriction of dislocation. In case consider that x ≈ d (several units of b), than length of dislocation segments, participating in the thermally activated process, must be also several units of b.

Conclusions The stress strain curves of Mg-x-Li3Al alloys are strongly dependent on the temperature. The flow stresses are decreasing with increasing temperature. The difference of the maximum stress and the yield stress strongly decreases with increasing temperature. At the temperatures about 200 °C the difference is very low, or almost none for Mg8Li3Al and Mg12Li3Al. The activation volume as a characteristics of thermally activated mechanism decreases with increasing stress.

127

Acknowledgements The authors are grateful for the support offered by the Grant Agency of the Czech Republic (Grant No. 106/03/0843).

References [1] H. Haferkamp, C. Jaschik, P. Juchmann, V. Kaese, M. Niemayer and P. Tai, Material Wissenschaft und Werkstofftechnik 2000, 32, p. 25. [2] K. Mathis, Z. Trojanová, P. Lukáč, Mater. Sci. Engn. 2002, A 327, p. 171. [3] V. I. Dotsentko, Phys Stat. Sol. 1979 (b), 93, p. 11. [4] A. R. Rosenfield, G. T. Hahn, A. L. Bement Jr., R. I. Jaffer, Dislocation Dynamics, McGraw-Hill Book Company, New York, USA, 1968.

The Effect of Grain Size on the Bolt Load Retention Behaviour of AMC-SC1 C.J.Bettles1, Lihui Zheng2, K.Venkatesan1 and Ma Qian2 1

CRC for Cast Metals Manufacturing (CAST), CSIRO Manufacturing & Infrastructure Technology, Melbourne, Australia. 2 CRC for Cast Metals Manufacturing (CAST), The University of Queensland, St Lucia, Australia

Abstract The alloy developed by the CRC for Cast Metals Manufacturing in collaboration with the Australian Magnesium Corporation, AMC-SC1, is a sand cast, heat treatable alloy, and its properties, therefore, may be dependent on the grain size achieved within the casting. One important property in relation to engine block applications is bolt load retention. It is shown that, within the range of grain sizes from 50 μm to 550 μm, the overall bolt load retention behaviour of AMC-SC1 at 423 – 450 K and loads of 8 kN and 11 kN is only mildly dependent on the absolute grain size. However, the individual factors contributing to the overall bolt load retention behaviour are themselves much more responsive to changes in grain size. For reliable BLR behaviour at 8 kN and 11 kN loadings, the grain size of AMC SC1 should be maintained below 200 μm.

Introduction It has been recognised for some time that magnesium alloys may be instrumental in achieving the weight reduction targets imposed on the automotive industry. There has been considerable interest in developing high pressure die cast magnesium alloys for engine applications [1, 2, 3], but very little has been reported on the development of elevated temperature alloys for precision casting processes. The Australian Magnesium Corporation, together with the CRC for Cast Metals Manufacturing (CAST), has developed a precision sand casting magnesium alloy, AMC-SC1, which can be used in conjunction with a novel crankcase design concept to produce a lightweight engine [4, 5]. The mechanical properties of such precision sand cast alloys are dependent on both the as-cast grain size and the subsequent heat treatment conditions. The precipitation sequence developed through an appropriate heat treatment is essentially independent of the size of the casting and the microstructure, and the contribution to the overall strength from precipitation hardening is reproducible and repeatable. The grain size, however, can vary from casting to casting, depending on the level of grain refinement in the melt and, to a lesser extent, the cooling rate. It is therefore important to understand the role that grain size plays in the final mechanical properties of an alloy and, on a practical note, to determine the maximum grain size that can be tolerated whilst maintaining the relevant property specification. One such property, which is extremely important in terms of the engine performance, is the bolt load retention (BLR) behaviour under the cyclic operating conditions of the en-

129 gine. The overall BLR behaviour is, in reality, a combination of two separate effects: Compressive yielding and compressive creep. In order to improve the BLR response of an alloy it is necessary to examine these contributions separately to ensure that the appropriate behaviour is targeted for improvement. In this paper, we describe firstly the effect of grain size on the overall bolt load retention behaviour and secondly the individual contributions from both yielding and compressive creep.

Experimental Details The preparation of the plates of AMC-SC1 with different grain sizes, together with the microstructures and mechanical properties achieved, has been described in another paper presented at this conference [6]. By varying the level of grain refinement (through the addition of varying levels of AM-cast [7]) and applying a chill to the mould, grain sizes ranging from a low of 46 μm to a high of 536 μm were achieved. The BLR behaviour was determined using an assembly similar to that described by Pettersen and Fairchild [1]. The two identical alloy bosses used in the assembly were each 16mm outside diameter, 8.3 mm inside diameter and 15 mm high. M8 high tensile bolts equipped with a load cell were used to apply the load to the assembly. Figure 1 shows a view of a fully bolted assembly used in this work. The test conditions chosen for this evaluation were 423 K and 450 K, with applied loads of 8 kN and 11 kN (corresponding to compressive stresses of 54.4 MPa and 74.8 MPa respectively). These loads were applied at room temperature, prior to the assemblies being placed in an oven at the test temperature. The duration of the test was restricted to 100 hrs, after which the assemblies were allowed to cool in air overnight to room temperature. The load on the bosses was recorded continuously using a data logger throughout the duration of the test.

Alloy bosses Load cell

Figure 1. A view of a typical fully bolted assembly .

Results and Discussion A generalised BLR relaxation curve is shown in Figure 2, together with the critical loads which define the BLR behaviour. These four loads are [8]: PI, the initial load applied to the assembly at room temperature (8kN or 11kN in this series of tests), PK, the maximum load reached in the assembly during heating to the test temperature,

130 PR, the load retained at the test temperature after 100 hours (immediately prior to returning to room temperature) and PF, the final load retained by the assembly once it has been returned to room temperature. The initial increase in load observed, (PK – PI), is due to thermal expansion mismatch effects between the alloy of the bosses and the bolt. The assembly may not have reached the operating temperature when PK is achieved, as the yielding behaviour of the alloy will determine the maximum load that can be sustained. 14

PK

12

PR

Bolt Load (kN)

10 8

PI

PF

6 4 2 0 0

20

40

60

80 100 Time (h)

120

140

160

18

Figure 2. Generalised bolt load retention curve showing the critical loads measured during a standard BLR test. Table 1. Critical loads on the relaxation curves for each grain size under the four different testing conditions.

Condition Grain size PI (μm) 423K 8kN

423K 11kN

450K 8kN

450K 11kN

536 337 235 110 46 536 337 235 110 46 536 337 235 110 46 536 337 235 110 46

8.10 7.97 7.97 7.92 8.27 11.00 11.39 12.00 11.44 11.20 8.31 8.84 8.39 8.72 8.20 11.18 11.50 11.50 11.55 11.19

PK

PR

PF

PF/PI PR/PK PI-PF Creep Yield

8.78 8.80 9.59 9.17 10.02 11.85 12.00 12.88 12.63 12.55 9.52 10.62 11.02 10.11 11.00 11.63 12.48 12.58 12.68 12.51

8.59 8.66 9.16 8.75 9.84 11.40 11.56 12.58 12.36 12.27 9.45 9.99 10.28 9.04 10.40 11.18 11.57 11.33 12.00 11.69

6.50 6.95 6.25 6.95 7.39 9.08 9.74 9.70 9.97 9.86 6.49 6.31 6.54 6.63 6.40 7.59 8.29 8.88 9.15 8.73

0.80 0.87 0.78 0.88 0.85 0.83 0.86 0.81 0.87 0.88 0.78 0.71 0.78 0.76 0.78 0.68 0.72 0.77 0.79 0.78

0.98 0.98 0.95 0.95 0.96 0.96 0.96 0.98 0.98 0.98 0.99 0.94 0.93 0.89 0.95 0.96 0.93 0.90 0.95 0.93

1.60 1.02 1.72 0.97 0.88 1.92 1.65 2.30 1.47 1.34 1.82 2.53 1.85 2.09 1.8 3.59 3.21 2.62 2.40 2.36

0.19 0.14 0.43 0.42 0.18 0.45 0.44 0.30 0.27 0.28 0.07 0.63 0.74 1.07 0.6 0.45 0.91 1.25 0.68 0.82

1.41 0.88 1.29 0.55 0.70 1.47 1.21 2.00 1.20 1.06 1.75 1.90 1.11 1.02 1.20 3.14 2.30 1.47 1.72 1.54

131 The ratios PF/PI and PR/PK give the fraction retained at room temperature and the test temperature respectively, and are the most commonly reported results from bolt load retention testing. The overall relaxation behaviour (PI – PF), however, is made up of two components – a compressive creep component (PK – PR), and a component due to yielding, which is obtained by difference. These important factors for the alloys of different grain size are given in Table 1. In terms of the overall BLR behaviour and its effect on engine performance, the important ratio is PF/PI, as the engine is cycled from room temperature to the operating temperature, and back to room temperature again. Figure 3 is a plot of the PF/PI data shown in Table 1. It is clear that at a load of 8 kN, at both 423 and 450 K, the ratio PF/PI is independent of grain size. By contrast, at the higher load of 11kN, there is a decline in the bolt load retained with an increase in grain size at both temperatures. The compressive yield strengths of the alloys at both temperatures are given in Table 2, together with the equivalent loads to reach these stresses in a standard BLR test. On the basis of this information alone, it may appear that all tests carried out at 8kN load are beneath the load limit for yielding, and that at 11kN load only the alloys with the two largest grain sizes would exceed the yield point. However, there is a significant increase in load as the assembly is heated to the test temperature, as shown in Table 1 (i.e., from PI to PK), and it is likely that during this heating phase (which lasts only a matter of minutes) more of the alloys will be subjected to loads in excess of the yield strength. In order to determine the effect of grain size on the individual components of the BLR response it is necessary, therefore, to take into account the value of PK, as this will have an effect on the subsequent relaxation behaviour. At 450 K the BLR response changes from being independent of grain size at an initial load of 8 kN to showing a linear decline with increasing grain size at 11 kN (Figure 3). The contributions from yield and creep at these two initial loads have been plotted in Figure 4, after being normalised by the appropriate value for PK. 1.00 0.95

1.00

(a)

0.95

0.90

423K

0.80 0.75

0.80 0.75

450K

0.70

423K

0.85 PF/PI

PF/PI

0.85

(b)

0.90

0.70

450K

0.65

0.65

0.60

0.60 0

200

400

Grain Siz e (µm)

600

0

200

400

600

Grain Siz e (µm)

Figure 3. The effect of grain size on the overall bolt load retention behaviour, as indicated by the ratio PF/PI, at (a) 8 kN and (b) 11 kN.

132 Table 2. Compressive yield strengths and equivalent boss loads at 423K and 450K for the various grain sizes of AMC-SC1.

Grain size (μm) 423K CYS (MPa) 536 60 337 73 235 81 110 90 56* 108

450K Equiv. Load (kN) CYS (MPa) 8.8 61 10.7 63 11.9 82 13.2 95 15.9 102

Equiv. Load (kN) 9.0 9.3 12.1 13.9 15.0

* the CYS measurements were carried out on samples from a different part of the casting where the grain size was slightly larger than the reported 46 μm.

Comparison of the PK values with the respective compressive yield strengths (CYS) at each grain size shows that at the nominal 8 kN initial load the PK values have exceeded the CYS at grain sizes of 337 μm and above, and at 11 kN the yield strength is exceeded at 235 μm and above. This is consistent with the observations in Figure 4(a), where the yield contribution to the overall behaviour is significant at the larger grain sizes but levels out below approximately 200 μm, irrespective of the initial applied load. At the smaller grain sizes, the yield contribution is relatively independent of the as-cast grain size. 0.300

0.120

(a)

0.250

0.100

k

Creep/P

Yield/P k

0.200 0.150 0.100 0.050

(b)

0.080 0.060 0.040 0.020

0.000 0

200

400

Grain Siz e

600

0.000 0

200

400

600

Grain Siz e

Figure 4. The effect of grain size on (a) the yield components and (b) the creep components of the overall bolt load retention behaviour at 450K, where the dashed lines are at 8kN and the solid lines are at 11kN load.

The creep contributions at 450 K, for both 8 kN and 11 kN initial loads, show a negative parabolic-type relationship, reaching a maximum at intermediate grain sizes. The results at 8 kN load are somewhat confusing, and the data for the 110 μm plate may be treated with some suspicion, as the unusually high creep contribution is not observed at the higher 11 kN initial loading. In general, at the larger grain sizes, the creep contribution is higher at the higher initial loading, but at the finest grain size the contributions are similarly low. This compressive creep behaviour mirrors the behaviour seen in the tensile creep tests on the same materials [6], where at grain sizes in excess of 200 μm the steady state creep rate decreased as the grain size increased, and at smaller grain sizes the creep rate decreased with decreasing grain size. It is believed that the fully grain refined microstructure of AMC-SC1, with intermetallic material at grain boundaries and triple points, reduces the effect of grain boundary sliding and the dominant mechanism is changed to one of dislocation climb.

133

Conclusions 1. The overall BLR behaviour of AMC-SC1 at 423 K, and 8 kN or 11 kN load, is independent of grain size, up to 536 μm. 2. At 450 K, the bolt load retained at 8 kN and 11 kN decreases as the grain size increases (up to 536 μm). 3. At 450 K, the yield contribution to the stress relaxation is stable at grain sizes less than 235 μm, and increases as the grain size increases to 536 μm. 4. At 450 K, the creep contribution to the stress relaxation is at a maximum at a grain size of 235 μm. 5. For reliable BLR behaviour at 8 kN and 11k N loadings, the grain size of AMC-SC1 should be maintained below 200 μm.

Acknowledgements The authors wish to thank Mr D. Graham of The University of Queensland for his assistance in the preparation and heat treatment of the cast plates. This work was conducted under an alliance between the Australian Magnesium Corporation and the Australian Cooperative Research Centre for Cast Metals Manufacturing. The CRC for Cast Metals Manufacturing (CAST) was established under and is supported by the Australian Government’s Cooperative Research Centres Scheme.

References [1] F. C. Chen, J. W. Jones, T. A. McGinn, J. E. Kearns, A. J. Nielsen and J. E. Allison, “Bolt-Load Retention and Creep of Die-Cast Magnesium Alloys” SAE Technical Paper 970325, 1997. [2] K. Y. Sohn, J. W. Jones, J. Berkmortel, H. Hu and J. E. Allison, “Creep and Bolt Load Retention Behaviour of Die Cast Magnesium Alloys for High Temperature Applications; Part 2 of 2” SAE Technical Paper 2000-01-1120, 2000. [3] I. P. Moreno, K. Y. Sohn, J. W. Jones and J. E. Allison, “Bolt-Load Retention Behaviour of a Die Cast Magnesium-Rare Earth Alloy” SAE Technical Paper 2001-010425, 2001. [4] C. J Bettles, C. T. Forwood, D. S. Jones, J. R. Griffiths, M. T. Frost, D. H. St. John, Ma Qian, G-L. Song and J. F. Nie, “AMC-SC1: A New magnesium Alloy Suitable for Powertrain Applications” SAE Technical Paper 2003-01-1365, 2003. [5] W. Schoeffmann, F. Beste, M. Atzwanger, H. Hick. and U. Sauerwein, “Lightweight Engine Structures – Mechanical, Acoustic and Production Aspects” SAE Technical Paper 2003-01-0406, 2003. [6] G. Dunlop, C. J. Bettles, J. R. Griffiths, K. Venkatesan, Li Zheng and Ma Qian, “The Effect of Grain Size on the Mechanical Properties of AMC-SC1” International Conference on Magnesium Alloys and Their Applications, 2003. [7] Ma Qian, D. H. St.John and M. T. Frost, International Conference on Magnesium Alloys and Their Applications, 2003, Wolfsburg. [8] K. Pettersen and S. Fairchild, “Stress relaxation in Bolted Joints of Die Cast Magnesium Components” SAE Technical Paper 970326, 1997.

High Temperature Behaviour of a new Mg-Ca-Zn Alloy A. Zilberov, G. R. Goren-Muginstein and M. Bamberger Department of Material Eng., Technion IIT, Haifa 32000, Israel

1

Introduction

In recent years, light Mg alloys have become a center of interest among engineers and scientist. The combination of lightweight, good strength and castability makes Mg alloys a promising engineering material of the automotive and aviation industries. Most common and commercial Mg alloys contain Al as a major alloying element. Al improves the castability and mechanical properties at room temperature, however at temperatures above 120 °C, Mg-Al bases-alloys, like AZ91, soften due to the existence of the γ-Al12Mg17 phase. At elevated temperatures, the Al solution hardening decreases and the γ-phase precipitates softens [1, 2]. The solution is to either add additional alloying elements that will reduce the amount of γ-phase by replacing some or all the Al with Ca and Zn [3–5]. Ca also behaves like as a grain refiner [6], and creates the stable intermetallic Mg2Ca [4, 7]. The second additive element is Zn, which can create various Mg-Zn [8] compounds and together with Ca creates Mg6Ca3Zn2 [4]. The goal of this research is to study the precipitation process at a predefined Mg-Zn-Ca alloy, at the as cast condition in order to investigate its structural stability at elevated temperatures. In addition, the alloy was solution treated in order to achieve a better understanding of the precipitation mechanisms and to test the alloy’s response to precipitation hardening.

2

Experimental

The alloy was prepared from pure a Mg, Zn and Mg-30%wt Ca pre-alloy, the raw materials were melted in cemented graphite crucible under Ar atmosphere. The melt was kept at 700 °C to ensure dissolution of alloying elements and then cast into a cold AISI 1020 steel mold, forming disks 9 mm thick and 60 mm in diameter. The samples for aging were encapsulated in a quartz tube filled with Ar, at a partial pressure of 397 mm Hg, for solution treatment. In order to eliminate the local remelting of the low-melting eutectic, the solution treatment was carried out in two steps; a) the alloys were held at 300 °C for 48 h, b) then heated to the holding temperature of 465 °C at 1 °C/hr and kept there for additional 96 h, followed by quenching in water. The solution treated samples were sliced and aged in a melted mixture of 50% sodium nitrate and 50% potassium nitrate. The treatment was preformed at different temperatures (150 °C–225 °C) for various times (1–96 h) followed by water quenching. Thermal exposure of the as-cast samples was done in the same manner as the aging treatments, but for longer times ranging betweSen 8 and 768 hrs. All samples were cut and polished for SEM, HRSEM, EDS, Auger spectroscopy, and XRD inspection.

135

3

Results and discussion

3.1

The investigated alloy

In order to find the most suitable composition for this research, a phase diagram simulation was made utilizing the Thermo-Calc program. The iso-thermal section at 465 °C in the ternary phase diagram of Mg-Zn-Ca is shown in fig. 1. In order to avoid grain boundary melting during solution treatments the following composition was chosen; Mg1.6 wt% Ca-3.2 wt% Zn.

Figure 1: Calculated Iso-Thermal cross-section at 465 °C in the Mg-Ca-Zn system

The calculated isotherm predicts the existence of at list two phases; α-Mg matrix and Mg2Ca. In order to design the solution treatments a second Thermo-Calc calculation was made. Fig. 2 presents the calculated phase fraction vs. the temperature for Mg-1.6 wt% Ca-3.2 wt% Zn. As one can see the liquid transforms into a eutectic structure with a third phase Mg6Zn3Ca2.

Figure 2: Calculated phase fraction vs. temperature for Mg-1.6 wt% Ca-3.2 wt% Zn

The as-cast structure of this alloy is shown in fig. 3. The alloy contains three phases; αMg matrix, Mg2Ca in the grain boundaries and Mg6Zn3Ca2 in eutectic structure – see also table 1, as calculated by the Thermo-Calc software. XRD spectra for the as-cast alloy shows the same phases.

136

Figure 3: SEM micrograph of the as-cast structure Table 1: Composition of the various phases in the as-cast state

EDS Auger spectroscopy Phases 3.2

α-Mg matrix [%at]

Gray phase [%at ]

White phase [%at]

Mg-0.3Ca-0.7Zn Mg-0.3Ca-0.4Zn α-Mg matrix

Mg-26Ca-2Zn Mg-29.3Ca-2.9Zn Mg2Ca

Mg-7.2Ca-13Zn Mg-18.6Ca-24Zn Mg6Zn3Ca2

The precipitation process for solution treated and aged samples

Fig. 4 shows the variation in hardness with the aging time at various temperatures. The highest hardness of 74 HV was observed after 8 hrs at 175 °C. The decrease of hardness at that temperature in the following 96 hrs is relatively minor. At 200 °C and 150 °C it seems that there are tow maximum picks. It is assumed that there’s a second pick at 150 °C but further investigation is required in order to verify it.

Figure 4: Hardness curves after several aging temperature as a function of time

137 Fig. 5a and b show SEM micrographs of the solution treated and aged samples correspondinglye. After solution treatment the grain boundary phase Mg6Zn3Ca2 was almost completely dissolved in the matrix, but the Mg2Ca coarsened and spheroidized. On the edge of each gray Mg2Ca grain a white trace of Mg6Zn3Ca2 is discernible. After aging for 96 hrs. no precipitation can be seen in this magnification.

Figure 5: SEM micrographs of: a) solution treated alloy and b) aged for 96 hrs at 225 °C

HRSEM micrographs of the aged sample are presented in fig. 6a and b. Round and oval precipitates of about 100–200 nm detected.

Figure 6: HRSEM micrographs of the aged alloy

3.3

Structural stability of the as-cast alloy at elevated temperature

The as-cast alloy was exposed to various temperatures for up to 32 days. Fig. 7 shows the hardness variation with time. The hardness in those samples is almost constant, hence this alloy makes a very promising candidate for service at elevated temperatures. Closer look at the microstructure evolution during the thermal exposure after 32 days at 200 °C (fig. 8) reveals small precipitates near the grain boundaries and at triple junctions. After casting, the area near the grain boundaries is enriched with the alloying elements, above their equilibrium composition, and when exposed to elevated temperatures, an aging process takes place. In addition coarsening of the grain boundaries is evident [5]. As the alloy

138 solution hardening and grain boundaries hardening decline with time, precipitation hardening compensates for it; hence the hardness is relatively constant at 68 ± 5 HV.

Figure 7: Hardness variation of the thermally exposed as cast samples for 32 days at 200 °C

Figure 8: SEM micrographs of the thermally exposed as cast samples after 32 days at 200 °C, etched by nitric acid

EDS examination of the precipitants seen in Fig. 8b reveals that these are MgZn. This phase was not found by the Thermo-Calc simulation presented in fig. 2. Both the calculations and the EDS finding need to be re-examined. XRD spectra of those samples did not reveal the existence of a forth phase in those samples.

4

Summary and conclusions

•

Mg-1.6 wt% Ca-3.2 wt% Zn was found to be suitable for both aging treatments and working at elevated temperatures. The as-cast alloy consists of three major phases; α-Mg matrix, Mg2Ca in the grain boundaries and Mg6Zn3Ca2 in eutectic structure.

•

139 • • • • •

Mg2Ca dose not dissolve during solution treatment. Maximum hardness was observed after 8 hrs at 175 °C, 74 HV. During aging and exposure to high temperature precipitates MgZn. Coarsening of the grain boundaries takes place during thermal exposure. The hardness is almost constant during the thermal exposure, hence the alloy makes a promising candidate for service at elevated temperatures.

5

Bibliography

[1] M. Regev, A. Rosen, and M. Bamberger, Met. Trans., 2001, 32A, pp. 1335–1345 [2] W. Blum, B. Watzinger, and P. Weidinger, Proc. Of the Mg Alloys and their Application Conf. (B. L. Mordike and K. U. Kainer eds) Werksoft Informationsgesellschaft mbH, Wolfsburg, Germany, 1998, pp. 49–60 [3] A. A. Lou and B. R. Powel in Mg Technology 2001 (J. Hryn Editor) Proc. Of the TMS symposium, 2001, pp. 137–144 [4] J. F. Nie and B. C. Muddle, Scripta Mater., 1997, 37, pp. 1475–1481 [5] A. Finkel, L. Shepeleva, M. Bamberger and E. Rabkin, in Mg Technology 2002, (H. I. Kaplan Editor) Proc. Of the TMS symposium, 2002, pp. 81–86 [6] Y. Guangyin, L. Manping, D. Wenjiang, and A. Inoue, Mater. Sci. Eng, 2003, 00A, pp. 1–7. [7] T. Miyazaki, J. Kaneko and M. Sugamata, Mare. Sci. Eng, 1994, A181/182, pp. 1410–1420 [8] A. A. Nayeb-Hashemi and J. B. Clark, in ASM International, 1998

Investigation of die Cast Mg-Al Alloys with Additions of Sn, Si and Sb Dr. Per Bakke, Dr. Ketil Pettersen, Dr. Håkon Westengen Hydro Aluminium a.s., Magnesium Competence Centre, Porsgrunn, Norway

1

Introduction

Magnesium die casting alloys based on Al as the main alloying element can broadly be grouped into two categories; those where the third element(s) form stable phases with Al, and those where the third element(s) form precipitates with Mg [1]. This is illustrated in Fig. 1. Alloys with Ca (AX), Sr (AJ) and RE (AE) belong to the first category, while alloys with Si (AS), Sb and Sn are examples from the second. Only few examples are provided where alloying elements from the different categories are combined. The reason for this is the free energy of formation (ΔG) of intermetallic compounds is closely related to the groups of the periodic table of elements, leading to a low mutual Figure 1. Elements forming stable intermetallic phases in solubility of the elements from the two the Mg-Al base system. groups. 1.1

History of Mg alloys with Sn, Sb and Si

Magnesium alloys with Sn additions was first mentioned in a 1934 patent [2]. According to Emley [3] Sn in Mg alloys confers some strengthening at the expense of ductility. Recent Mg alloy patents from Japan [4-6] include Sn among the alloying elements, beneficial for strength and creep properties. This indicates that Sn may be an interesting element in die cast Mg alloys where high strength and creep resistance are important, although corrosion may be a problem especially in alloys with Al and/or Zn [7,8]. Tin was a constituent of some early Mg based hammer-forging alloys, especially an alloy with 5% Al and 3% Sn, [8]. However, as far as we know, there are no Mg alloys with Sn in use today. Magnesium alloys with Sb additions were first patented back in 1936 [9]; a sand casting alloy for engines where Sb was added for improving the hardness. In the 1970’s [10-12] it was discovered that small amounts of Sb added to AS alloys improve the corrosion properties. Recently, work on microadditions of Sb (<1 wt%) to AZ91 concluded that for permanent mould cast and T4 treated specimen a significant increase in yield strength and creep resistance up to 200 °C was obtained at the expense of ductility [13, 14]. Alloys with Sb are not in commercial use today.

141 Die cast magnesium alloys with silicon have been used commercially for a long time [11, 14-21]. The VW beetle engine for example, was first produced in AZ81, but later changed to AS41 and AS21 in the early 1970’s due to significantly improved creep properties at 150 °C [20]. Due to inadequate corrosion properties for some applications, AS21 has recently been modified with minor additions of RE [22-24]. However, as discussed earlier, due to restricted mutual solubility the amount of RE that can dissolve in the AS alloys is too small to have any effect on any properties but corrosion. 1.2

Thermodynamic data

Some thermodynamic data for the Mg-Si, Mg-Sn and Mg-Sb systems are shown in Table 1. Like Mg2Si in Mg-Al-Si alloys, Mg2Sn in Mg-Al-Sn alloys and Mg3Sb2 in Mg-Al-Sb alloys can be expected to facilitate a stabilised matrix structure which is a necessity for a creep resistant alloy [1, 25]. Due to a small negative enthalpy of mixing, Sn dissolves into liquid Mg in an exothermic reaction [26]. This combined with the large solubility of Sn in liquid Mg and the low melting point of Sn (232 °C) makes Sn very easy to alloy, and 100% recovery can be expected. Antimony has an even larger negative enthalpy of mixing and a significant temperature increase can be experienced during alloying. Pure Sb melts at 630.8 °C. Silicon melts at 1410 °C, hence the dissolution process is mass transfer controlled. As a result, efficient alloying of Si requires powerful stirring to continuously keep the solid Si in contact with molten alloy of uniform composition, minimising compositional gradients [27]. Table 1. Some thermodynamic data for the Mg-Si, Mg-Sn and Mg-Sb systems [28].

System

Eutectic comp. Solid solubility Eutectic temp. Secondary Melting [at%/wt%] [at%/wt%] [°C] phase temp. [°C]

Mg-Si Mg-Sn Mg-Sb

1.16/1.34 10.7/36.9 11.5/39.4

0.04/0.05 3.35/14.48 negligible

637.6 561.2 629

Mg2Si Mg2Sn Mg3Sb2

1085 770.5 1245

From Table 1 it is seen that there is practically no solid solubility of Sb in Mg, and Sb additions alone do not provide any solid solution strengthening to the alloy [1, 25]. The solid solubility of Si in Mg is also too low to have any significant impact. Nevertheless, belonging to group VA, Sb and Si form extremely stable phases with the elements of group IIA (Mg, Ca, Sr,…), as indicated in Fig. 1. On the other hand, Sn has considerable solid solubility in Mg, and some solid solution strengthening may be expected. Also, the stable Mg2Sn particles can be anticipated to have a positive influence on the creep properties. In AS-alloys the optimum Al-content regarding creep is found to be around 2 wt% [24] Another limitation for further improvement of the creep resistance of AS21 is the limited liquid solubility of Si in the Mg-Al base alloy, effectively restricting Si contents to roughly 2 wt%. An approach along this line to completely or partly replace Si with Sb in the AS alloy has recently been reported [29]. With the similarities (and differences) between Sn and Si in mind, die cast Mg-Al-Sn(-Mn) alloys have been investigated, extending from binary Mg-Sn alloys. Ternary Mg-Al-Sb and Mg-Al–Si alloys, including hypereutectic alloys with up to 3% Si, have been examined. The Mg-Al-Si-Sn system has also been looked into. Testing included tensile properties at room temperature and creep properties at 150 °C/50 MPa. Corrosion issues were not examined.

142

2

Experimental

Die casting was in part performed on the Frech 200 ton cold chamber machine, using the 3-cavity test bar die “TBD-e” (Fig. 2), and in part on the Buhler Evolution, 420 ton cold chamber machine with a combined plate/bar die (Fig. 3). The metal was hand ladled. A metal temperature of 700 °C was used during die casting. However, for some of the higher alloyed compositions the melt temperature was reduced to 690 °C. The oil heater temperature was 180 °C, and Klüber HP1-415 lubricant was applied by hand sprayFigure 2. The test bar die, TBD-e for the 200 ton Frech cold ing. For each composition a minimum of 15 good samples were cast. chamber machine. Standard 6 mm test bars from each composition were tensile tested at a 10 kN Instron machine at room in accordance with ISO 6892. In all tests a strain rate of 1.5 mm/min was used until 0.5% strain, whereafter the strain rate was increased to 10 mm/min. Two specimens from a series of alloys were tensile creep tested for minimum 150 hrs at 150 °C/50 MPa, in accordance with ASTM E139. Chemical analyses were done on permanently mould cast medallions using XRF-Uniquant, ICP-AES and FFigure 3. The Plate/bar die (version 0) for the Bühler EvoluAAS. tion 420 ton machine.

3

Results and discussions

3.1

Tensile testing at ambient temperature

The chemical analyses of the investigated alloys are shown in Table 2 together with tensile yield strength (TYS), ultimate yield strength (UTS) and elongation at fracture, Ef. The tensile data reported are the average over at least 10 specimens. Standard deviations are indicated. TYS and UTS at room temperature for all alloys is shown in Figs. 4-6. In Figs. 4 and 5 the results are grouped according to the Al content, 0, 1-2 wt%, 2-3%, 5-6% and 8-10%. Results of AZ91, AM50 and AS21X are included. These are obtained on specimens cast with the same die on the same HPDC machine; the AZ91 results are from one single test series, while the AS21X results are averaged values from a number of test series [24]. It is seen from Fig 4 that TYS increases apparently linearly with increasing content (at%) of alloying elements [1]. At low contents of alloying elements it seems that the binary Mg-Sn alloys constitute an upper barrier for the tensile yield strength. For example, the binary Mg-Sn alloy with 26 wt% Sn (6.8 at%) has TYS as high as 165 MPa. At very

143 high contents of alloying elements, here obtained by Sn and Al at 14 and 10wt%, respectively, tensile yield strength exceeding 200 MPa can be obtained. Also, the alloy with 7 wt% Sn and 8 wt% Al has TYS at 186 MPa. In comparison, TYS of AZ91 is at 150 MPa. Table 2: Chemical analyses and average TYS, UTS and Ef. In the table the ASTM design-ations for tin and silicon (T and S, respectively) are used. For antimony Sb is used.

Al 2.0 1.9 1.8 1.8 1.91 1.96 1.92 1.71 1.59

Mn 0.4 0.3 0.3 0.3 0.22 0.28 0.32 0.24 0.23

0.05 1.21 1.77 8.32 1.71 9.83 2.6 5.0 2.7 5.2

0.19

Si 1.76 2.27 2.64 2.00 1.94

0.21 0.26 0.20 0.33

Sb Sn 5.0 9.1 11.8 13.0 6.3 12.4 6.3 12.3 19.2 26.3 13.1 1.8 3.4 7.3 14.0 14.2 6.1 6.0 12.5 12.3

TYS [MPa] 114 ± 1.8 124 ± 1.4 123 ± 2.5 123 ± 3.0 116 ± 2.7 115 ± 6.0 112 ± 2.7 128 ± 2.1 129 ± 2.0 88.4 ± 3.7 111.8 ± 4.0 135.9 ± 4.3 165.4 ± 2.5 118.0 ± 3.6 96.6 ± 11.7 97.5 ± 2.3 186.9 ± 4.5 134.0 ± 2.5 209.0 ± 3.6 113.1 ± 2.5 134.0 ± 1.8 134.0 ± 1.5 157.2 ± 2.2 146 ± 6 120 90

250

300

200

250

UTS (MPa)

TYS (MPa)

Sample ID ASb25 ASb29 ASb2-12 ASb2-13 AS21 AS22 AS23 ASSb22-6 ASSb22-12 6 Sn 12 Sn 18 Sn 25 Sn TM12-0 AT12 AT23 AT10-6 AT2-12 AT10-12 ATM260 ATM560 ATM2-12-0 ATM5-12-0 AZ91 [30] AS21/AS21X [24] AM20 [31]

150

Mg-Sn(-Mn) Mg-Sn-Al, 1-2%Al) Mg-Sn-Al-Mn, 2-3% Al Mg-Sn-Al-Mn, 5-6% Al Mg-Sn-Al, 8-10% Al AZ91 AS21X

100 50 0 0

2

4

6

8

10

12

UTS [MPa] 184 ± 20 176 ± 11 173 ± 10 158 ± 10 205 ± 9 211 ± 10 213 ± 2 191 ± 9 174 ± 11 152.8 ± 14.3 168.1 ± 12.9 178.0 ± 7.1 202.4 ± 10.0 176.8 ± 17.1 211.8 ± 15.3 196.3 ± 23.1 244.8 ± 19.8 213.8 ± 12.4 247.8 ± 18.6 210.5 ± 27.1 238.3 ± 21.2 220.5 ± 25.6 224.4 ± 10.8 253 ± 8 230 200

200 150

Mg-Sn(-Mn) Mg-Sn-Al, 1-2%Al) Mg-Sn-Al-Mn, 2-3% Al Mg-Sn-Al-Mn, 5-6% Al Mg-Sn-Al, 8-10% Al AZ91 AS21X

100 50

14

at% alloying elements

Figure 4. Tensile yield strength vs. at% alloying elements in the Mg-Sn-Al(-Mn) system. AZ91 and AS21X are included for comparison. All specimens were cast in the TBD-e die (Fig. 2).

Ef [%] 6.9 ± 2.6 4.3 ± 1.0 3.9 ± 1.0 2.5 ± 0.9 7.2 ± 1.3 8.0 ± 1.4 7.6 ± 0.6 2.6 ± 0.6 1.4 ± 0.4 5.7 ± 1.4 4.5 ± 1.3 2.6 ± 0.8 1.7 ± 0.9 4.8 ± 1.6 14.7 ± 2.0 11.9 ± 4.7 1.9 ± 0.5 7.0 ± 1.9 1.1 ± 0.2 10.8 ± 4.8 9.3 ± 2.5 7.7 ± 2.5 4.2 ± 1.1 7.7 ± 0.9 12 21

0 0

2

4

6

8

10

12

14

at% alloying elements

Figure 5. Ultimate Tensile strength vs. at% alloying elements in the Mg-Sn-Al(-Mn) system. AZ91 and AS21X are included for comparison. All specimens were cast in the TBD-e die (Fig. 2).

144 250

16

Elongation (%)

25

Mg-Sn(-Mn) Mg-Sn-Al, 1-2%Al) Mg-Sn-Al-Mn, 2-3% Al Mg-Sn-Al-Mn, 5-6% Al Mg-Sn-Al, 8-10% Al AZ91 AS21X

14 12 10 8

Sb 20 Elongation (%)

UTS or TYS (MPa)

From Figs. 4 and 5 we see that the yield strength of the ASb alloys are in the same range as the AS alloys. As discussed earlier [1] TYS as 150 a function of alloying elements roughly follows a UTS Sb 100 TYS - Sb straight line, and the present data, including both UTS - Si TYS - Si 50 UTS - Si+Sb ASb, AS and ASSb, do not represent any excepTYS - Si+Sb 0 tions. 0 1 2 3 4 5 The results indicate that UTS and ductility of at% alloying elements the ASb alloys are lower than of AS alloys with Figure 6. UTS and TYS as a function of at% the same content (at%) of alloying elements. alloying elements (Sb, Si or both) in alloys with 2 wt% Al. The zero level represents AM20. Additions of 6-12 wt% Sb to AS22 lead to a Filled markers represent specimens cast in the significant drop in ductility, and also a slight TBD-e die (Fig.2) while open markers represent decrease in UTS. the plate/bar die (Fig. 3). Elongation as a function of at% alloying elements is shown in Figs. 7 and 8. It is seen that Al added to the Mg-Sn alloy significantly improves the ductility (Fig. 7). Antimony added to Mg-Al-Si alloys has a marked negative effect on the ductility (Fig. 8). 200

6 4

Si Si+Sb

15 10 5

2

0

0 0

2

4

6

8

10

12

14

at% alloying elements

Figure 7. Elongation vs. at% alloying elements in the Mg-Sn-Al(-Mn) system. Typical values of AZ91 and AS21X are included for comparison. All specimens were cast in the TBD-e die (Fig. 2).

3.2

0

1

2

3

4

5

at% alloying elements

Figure 8. Elongation at fracture as a function of at% alloying elements (Sb and/or Si) in alloys with 2 wt% Al. The zero level indicates AM20. Filled markers show specimens cast in the TBD-e die (Fig.2). Open markers represent the plate/bar die (Fig. 3).

Tensile creep

Creep curves for some of the alloys are shown in Figs. 9-13. Results of AM20, AZ91 and AS21X are included. Each of the plotted curves represent average of 2 - 5 individual test curves. It is seen that the creep performance of the binary Mg-Sn alloys improves strongly with increasing Sn content, Fig. 9. The alloy with 6 wt% Sn fairly quickly enters the phase III creep mode, while the alloy with 18% Sn has approximately identical creep curve with AS21X. The alloys with 12 and 25% Sn seem to undergo a significant creep during phase I, while their phase II creep rates are fairly similar to AS21X. It is not clear whether this creep during phase I is real or due to experimental errors. From Figs. 10 and 11 it can be seen that increasing Sn content improves the creep properties at 150 °C/50 MPa. At constant Sn level, increasing Al content is negative for the creep properties. From Fig. 12 it is obvious that the addition of Sb has a significant effect on reducing the creep strain relative to AM20. Increasing the Sb-content from 5% to the higher levels of 9 - 13 % give alloys with slightly improved creep strain compared to AS21. From Fig. 13 it is seen that increasing the Si content to hypereutectic levels in AS alloys has little effect on o the creep strain at 150 C/50 MPa. However, additions of Sb give a noticeable improvement.

145 10

10

AT12 AT23 AZ91 AT10-6 AT10-12

1

Mg 12Sn Mg 25Sn AS21X Mg 18Sn

0.1

Creep strain (%)

Creep strain (%)

Mg 6Sn

0.01

1

AT2-12

0.1

0.01 0

25

50

100

75

125

150

0

25

50

Time (hrs)

75

100

125

Figure 9. Creep strain curves for binary Mg-Sn o alloys obtained at 150 C/50 MPa. AS21X plotted as reference.

Figure 10. Creep strain curves for Mg-Al-Sn alloys obtained at 150 oC/50 MPa. AZ91 plotted as reference.

10

0.003

AM20

ASb25

ATM560 ATM260 ATM5-12-0 ATM2-12-0 TM12-0 AS21X

0.1

Creep strain (%)

Creep strain (%)

0.0025

1

150

Time (hrs)

ASb2-13

0.002

AS21

0.0015

ASb2-12 ASb29

0.001 0.0005

0.01 0

25

50

75

100

125

150

Time (hrs)

Figure 11. Creep strain curves for Mg-Al-Sn-Mn o alloys obtained at 150 C/50 MPa. AS21X plotted as reference. 0.003

0 0

25

50

75

100

125

150

Time (hrs)

Figure 12. Creep strain curves for Mg-Al-Sb alloys o obtained at 150 C/50 MPa. AM20 and AS21 plotted as reference.

Average creep rates at 100 hrs are shown in Figs 14-15. Also here the results are plotted as a function of at% 0.002 AS23 alloying elements. From Fig. 14 it is AS21 0.0015 AS22 seen that the creep rates of the alloys ASSb226 with 12 wt% Sn (with and without Mn) 0.001 are at the level of AS21X. All other 0.0005 alloys have higher creep rates. It is ob0 served that the minimum creep rates are 0 25 50 75 100 125 150 obtained around 12 wt% Sn, and increaTime (hrs) sing Al-content decreases the creep Figure 13. Creep strain curves for Mg-Al-Si and Mg-Alresistance. Addition of 0.2% Mn to Si-Sb alloys obtained at 150 C/50 MPa. AM20 and AS21 alloys with a significant amount of alplotted as reference loying elements does not have any significant impact on the creep behaviour. However, when Mn is added to Mg 2%Al to make AM20, a 10 times improvement in the creep rate is realised. AM20

Creep strain (%)

0.0025

o

146 1.00E-03

1.00E-03

1.00E-04

1.00E-04

1.00E-05

Mg-Sn(-Mn) Mg-Sn-Al, 1-2%Al) Mg-Sn-Al-Mn, 2-3% Al Mg-Sn-Al-Mn, 5-6% Al Mg-Sn-Al, 8-10% Al AZ91 AS21X

1.00E-06

1.00E-07 0

2

4

6

8

10

12

14

at% alloying elements

Figure 14. Average creep rates at 100 hrs as a function of at% alloying element (combinations of Sn, Al and/or Mn). All specimens were cast in the TBD-e die (Fig. 2)

Creep rate (hr -1)

-1

Creep rate (hr )

Sb Si Si+Sb AM20

1.00E-05 1.00E-06 1.00E-07 0

1

2

3

4

5

at% alloying elements

Figure 15. Average creep rates at 100 hrs as a function of at% alloying element (Sb, Si, Mn) in alloys with 2 wt% Al. The zero level indicates Mg 2Al. Filled markers show specimens cast in the TBD-e die (Fig.2). Open markers represent the plate/bar die (Fig. 3).

It is seen that AS21 (1.0 wt% Si) roughly corresponds to 5.0 wt% Sb, both with respect to creep rate and composition at an at% level. Slight improvements may be obtained by increasing the Sb content to 2 at% (10 wt%) or higher. It is also seen that the creep rate can be improved slightly by increasing the Si content to higher contents than in AS21. Furthermore, the creep rates as functions of at% Si or Sb are almost identical. The addition of Sb to AS22 has a marked positive effect on the creep rates. Notice that the sensitivity of the tensile creep test is at a level of 5 ⋅ 10-7 hr-1. Average creep rates at 100 hrs (average between 80 and 120 hrs) are used as estimates for minimum creep rates during phase II creep. This is true as long as the onset of phase III creep occur at Ⰷ150 hrs. Also, phase I creep should be terminated at Ⰶ100 hrs. Otherwise, the average creep rate at 100 hrs will slightly overestimate the minimum creep rate. 3.3

Metallography

Figure 16 shows the die cast microstructure of four of the alloys, TM12-0 (0.19% Mn, 13.1% Sn), ATM2-12-0 (0.20% Mn, 12.5% Sn, 2.7% Al), ATM5-12-0 (0.33% Mn, 12.3% Sn, 5.2% Al) and AT10-12 (14.2% Sn, 9.8% Al). It is seen that the Mg-Sn alloys with high Sn-contents have a more or less continuous grain boundary eutectic. With increasing additions of Al, the Mg-Sn eutectic phase is pushed together in coarser pockets at the grain nodes. Thus, the effect of grain boundary pinning is strongly reduced with increasing Al additions. In spite of Sn and Si belong to the same group of elements (IVA) and form intermetallic particles with Mg having the face-centered cubic CaF2 structure [32], Mg2Sn particles appear very differently from Mg2Si in AS alloys, which are characterised by the “Chinese script” structure [33]. Typical die cast structures of the ASb alloys are shown in Fig. 17. There is a large quantity of intermetallic particles of the type Mg3Sb2. In the ASb29 alloy the intermetallic particles are of a medium size and seem to be plate-shaped. The alloys ASb2-12 and ASb2-13 seem to have the same distribution of medium sized particles, but in addition there are significant numbers of coarse particles, as seen in Fig. 17. Additions of Sb beyond 9 wt% basically add coarse particles which have limited positive effects.

147 The SEM analyses of specimens from alloys TM12-2 (0.19% Mn, 13.1% Sn, Fig. 18, left) and AT10-6 (7.3% Sn, 8.3% Al, Fig. 18, right) show that the grain interior consists of a solid solution of both Al and Sn (spots L2, L3 and R1), with segre-gation towards the grain boundaries (R2). The intermetallic phases at the grain boundaries are Snrich, most probably Mg2Sn Figure 16. Optical micrographs showing the die cast microstructure of four (L1, L4, L5 and R5). With of the alloys, TM12-0 (0.19% Mn, 13.1% Sn, upper left), ATM2-12-0 high Al-content also an (0.20% Mn, 12.5% Sn, 2.7% Al, upper right), ATM5-12-0 (0.33% Mn, Al-rich phase, possibly Sn12.3% Sn, 5.2% Al, lower left) and AT10-12 (14.2% Sn, 9.8% Al, lower modified Mg Al is found 17 12 right). (R3 and R4). The solid solution of Sn in the matrix is expected to have a beneficial effect on the creep resistance. In combination with Al, however, this seems not to be the case. The most positive effect of Sn on creep, is with the compositions which form a continuous grain boundary eutectic in combination with a matrix with some solid solution. The creep resistant alFigure 17. Die cast structures (200x), ASb25 (upper left), ASb29 (upper loys within the Mg-Al right), ASb2-12 (lower left), and ASb2-13 (lower right). base system obtain their creep resistance by a relatively low content of Al, and addition of elements that form stable intermetallic phases within the grain mantle [1]. As indicated in Table 1, elements from groups IIA (Ca, Sr) and IIIA (RE) form very stable intermetallic particles with elements from groups IVA (Si) and VA (Sb). The mutual liquid solubility of the elements within each group can be significant, while it generally is limited between the elements from different groups. When using Si or Sb for particle formation in the alloy there is almost nothing to gain from adding for instance Ca, Sr or RE in addition, unless the temperature is raised significantly. With AS and ASb alloys, the options for potential improvement are thus in practice limited to Ge, Sn, Pb and Bi. The morphology and distribution of the phases vary significantly between the various alloying elements, and the effect on the mechanical properties varies to some extent accordingly.

148

Spot no. L1 L2 L3 L4 L5 Alloy

Al 0.22 0.00 0.00 0.23 0.05 0.05

Mn 0.35 0.06 0.16 0.04 0.17 0.09

Sn 11.44 0.98 1.18 10.50 9.00 3.00

Fe 0.44 0.00 0.00 0.00 0.00 0.00

Mg 87.54 98.95 98.66 89.23 90.78 96.86

Spot no. R1 R2 R3 R4 R5 Alloy

Al 4.47 7.98 19.01 28.11 3.19 8.03

Mn 0.03 0.00 0.00 0.06 0.14 0.00

Sn 0.89 1.68 1.72 1.33 26.50 1.60

Mg 94.61 90.34 79.26 70.50 70.17 90.37

Figure 18. SEM micrographs with EDS analyses of phases in alloys TM12-0 (0.19% Mn, 13.1% Sn, left) and AT10-6 (7.3% Sn, 8.3% Al, right). All analyses in the table in at%.

The tests may indicate that addition of up to 10 wt% Sb can be an alternative to the use of Si for obtaining creep resistance, provided that the relatively low tensile strength is acceptable. It is also demonstrated that the same creep properties but slightly better tensile strength can be obtained by increasing the Si content up to a hypereutectic level of 3 wt%. However, alloys with that high Si content may need very strict temperature control in the die casting shop to avoid significant and possibly detrimental precipitation of Si containing intermetallic particles, with loss of Si as an inevitable result [29].

4

Conclusions

• Alloys with high Sn and Al contents display tensile yield strength values above that of AZ91D, along with an equivalent UTS. The UTS and ductility of the ASb alloys are lower than what is observed in AS alloys with the same content (at %) of alloying elements. With increasing content of Sb in Mg-Al alloys, the TYS increases, while the ductility decreases. However, both AS22 and AS23 both have better yield strength than the ASb alloys. Additions of 6-12 wt% Sb to AS22 lead to a significant drop in ductility, and a slight decrease in UTS. • The creep resistance increases with increasing Sn or Sb-content; increasing the Alcontent reduces the creep resistance. The addition of 5 - 13 wt% Sb to AM20 leads to creep rates slightly better than AS21. Increasing the Si content up to a hypereutectic level of nearly 3 wt% leads to slightly better minimum creep rates than AS21. • The microstructure of the Mg-Al-Sn alloys consists of an α-Mg matrix with Al and Sn in solid solution, and a grain boundary eutectic of Mg-Mg2Sn. With increasing Alcontent, the Mg2Sn phase is pushed into pockets at the grain nodes, hence their effect on grain boundary pinning is reduced. There is no solid solubility of Sb in α-Mg. Additions beyond 9% Sb mainly leads to the formation of coarse particles, which don’t have any further positive effect on the properties. • The castability of Mg-Sn and Mg-Sb alloys is improved by addition of Al.

149

References [1] K. Pettersen, P. Bakke, D. Albright, Magnesium Technology 2002, (Ed.: H.I. Kaplan), TMS-AIME, 2002. 171-176 [2] M.E. Brooks, US Patent 1,957,934, Patented 8-05-1934 [3] E.F. Emley, Principles of Magnesium Technology, Pergamon Press, London, 1966. p.232 [4] K. Aikawa, K. Taketani, US Patent 5,304,260, Patented 19-04-1994 [5] T. Horie, et al., Japanese Patent 07,003,374 A, Filed 16-06-1993 [6] K. Nakamura, et al., European Patent Application, EP 1,108,799 A2, Filed 4-12-2000 [7] A. Beck, The Technology of Magnesium and its Alloys, F.A. Hughes & Co., London, 1940. p.286 [8] C.S. Roberts, Magnesium and its Alloys, John Wiley & Sons, New York, 1960. p.181 [9] F. Christen, US Patent 2,065,170, Patented 22-12-1936. [10] G.S. Foerster, Paper no G-T75-112, Proc. 8th SDCE, 1975. [11] G.S. Foerster, et al., Die Casting Eng., Jan.-Feb. 1977. 12–18. rd [12] G.S. Foerster, 33 IMA Conf., 1976. 35-39. [13] G. Yuan, Y. Sun, W. Ding, Scripta Mater., 43, 2000. 1009-1013 [14] G. Yuan, Y. Lu, X. Zheng, Y. Zhu, W. Ding, Acta Met. Sinica, 37, (1) Jan. 2001. 23-28 [15] G.S. Foerster, Norwegian Patent No. 121,753, Filed 19-02-1968 [16] G.S. Foerster, French Patent No. 1,555,251, Filed 29-02-1968 [17] K.E. Nelson, Paper no. 13, Proc. 6th SDCE, 1971 [18] S.L. Couling, Metals Eng. Quarterly, May 1972. 7-13 [19] E.W. Balzereit, Metallwissenschaft und Technik, 26, March 1972. 245-250 [20] F. Hollrigl-Rosta, E.Just, J.Kohler, H.-J.Melzer, Metall, 34 J, H12, Dec. 1980. 11381141 [21] S.L. Couling, US. Patent No. 529.885, Filed 05-12-1974 [22] K. Pettersen, J.I.Skar, M.Videm, Intl. Patent Application WO 0102,614, Filed 25-101999 [23] K. Pettersen, H. Westengen, J.I. Skar, M. Videm, L.-Y. Wei, Magnesium Alloys and their Applications, (Ed. K.U.Kainer), Wiley-VCH, Weinheim, Germany, 2000. 29-34 [24] K. Pettersen, P. Bakke, J.I. Skar, C. Tian, D. Albright, NADCA 2001. 193-201 [25] H. Westengen, P. Bakke, PSTAM, Osaka, Japan, Jan. 2003 [26] O. Kubaschewski, C.B. Alcock, Metallurgical Thermochemistry, 5th Ed., Pergamon Press, Oxford, UK, 1979 [27] R.I.L. Guthrie, Electric Furnace Conf. Proc., 1977. 30-41 [28] A.A. Nayeb-Hashemi, J.B. Clark, Phase Diagrams of Binary Magnesium Alloys, ASM Intl., Metals Park, Ohio, 1988 [29] P. Bakke, K. Pettersen, Proc. Materials Week 2002, Werkstoff-Informationsges., 2002 [30] T. Aune, Internal work at Hydro Magnesium [31] Hydro Magnesium, Technical Information Brochure, Nov. 1996 [32] C.S. Roberts, Magnesium and its Alloys, John Wiley & Sons, New York, 1960. p.35 [33] A. Luo, M.O. Pekguleryuz, J. Mat. Sci., 29, 1994. 5259-5271

150

An Electron Microscopy Investigation on as-cast AZ91D Alloy Modified with Nitrogen A.A. Kaya*, O. Yucel**, D. Eliezer***, E. Aghion * TUBITAK-MRC, MCTRI, Kocaeli, Turkey ** Istanbul Technical University, Istanbul, Turkey *** Ben-Gurion University of The Negev Department of Materials Engineering Beer-Sheva, Israel **** MRI, Beer-Sheva, Israel

Abstract An attempt has been made to modify the microstructure of the most widely used magnesium alloy, AZ91D, via introducing nitrogen into the alloy in molten state. The microstructures of the samples before and after nitrogen addition were compared by using scanning and transmission electron microscopy examinations. It was observed that despite the near-zero solubility of nitrogen in magnesium the microstructure of the small ingots cast after purging the molted alloy with nitrogen showed considerable modifications as compared to the control sample cast before purging operation. Nitrogen addition appeared to lead to the elimination of the classical lamellar morphology from the microstructure. Furthermore, the divorced eutectic structure commonly observed in AZ91 was greatly promoted in the modified microstructure leading to the formation of much more α (a solid solution of Mg and Al) particles in grain boundary coarse β (Al17Mg12) as compared to the unmodified structure. Such modifications may be expected to bring about considerable changes in mechanical behavior, or create an opportunity to assess the individual effect of the lamellar structure on corrosion and mechanical behavior of AZ91 alloy.

1

Introduction

AZ91 alloy has been subjected to numerous investigations and therefore its mechanical properties and microstructure as die cast material has been well established (1-8). However, a similar degree of maturity of knowledge on the precise individual effects of the microstructural constituents in cast structures is not apparent. Therefore this study was undertaken in an attempt to modify the as-cast microstructure of AZ91 ingot for creating an opportunity towards the assessment of the individual role of a particular classical microstructural constituent namely the discontinuous precipitation product. The purpose of modification of ingot structure is also not unwarranted. The as-cast ingot structure is not of interest for a number of reasons. A parallelism in terms of solidification rate may be drawn between the as-cast ingot structure and other castings, e.g. sand castings. Furthermore, when considering the use of ingots as the starting material for shaping processes other than direct casting, such as semi-solid injection molding which requires chipping of solid materials, the need to understand the resultant structure of the

151 primary solidification process, i.e. that of ingot, becomes obvious. Furthermore, a successful chemical microstructural modification may be directly applicable to the shaping processes, such as diecasting, that are of higher interest industrially. Solidification in AZ91 was studied by Petterson et al. (9,10) who reported that the growth of dendritic stem took place in two directions, < 1120 > and < 2245 >. Aghion and Bromfin (11) investigated the microstructure and properties of various AZ91 ingots by light and scanning electron microscopy techniques. Clark studied Mg-9%Al alloy by TEM after aging (12). Crawley et al. (13,14) later complimented Clark’s report, and furthermore suggested crystallographic orientation relationships between the α matrix and β products in aged structures revealing that the plates form parallel to the basal planes of the hexagonal matrix. Previous studies (12-14) on the aging behaviour of die-cast AZ91 revealed that the precipitates of β phase form with various morphologies that includes discontinuous and continuous precipitation products with lamellar, and plate morphologies, respectively. It was also reported that aging process did not result in the formation of additional morphological features. A difference was noticed only in the precipitate density and the size, double aging resulting in a higher density of β plates with a higher aspect ratio. In a study on AZ91E, Bettles (15) made a distinction between the previously reported plate shaped β and what he called ‘lozange’ shaped continuous precipitation products, and also revealed that trace additions of copper inhibited the discontinuous precipitation. Assessment of the individual effects of particular types of precipitate morphologies on the properties of the material is difficult as all of them exist in the same material. Some information has been reported regarding the effect of discontinuous precipitation on corrosion properties (16). Generally this kind of cellular microstructures are considered as detrimental to the toughness of the material. When evaluating the effect of discontinuous precipitation on mechanical properties the nodule size should also be considered. Elimination of this type of precipitation reaction in Mg-based alloys would be desirable as it is also claimed to accelerate creep at high temperatures (17). Furthermore, absence of this constituent would presumably allow for more extensive continuous precipitation leading to better strengthening effect. Coexistence of a few morphologies of what is essentially β, Al12Mg17, in the same material raises questions that are worth considering. Coarse β particles at the grain boundaries are the product of a divorced eutectic reaction from the aluminum-enriched part of the liquid metal that solidifies the last in a casting (11). Unfortunately, little has been proposed towards explaining the relative conditions for the formation of the competing continuous and discontinuous precipitate morphologies. The continuous precipitation products in AZ91 are known to nucleate on dislocations and twins in grain interiors (12-14). Increased precipitation in pre-deformed samples strongly supports this idea (12). However, what leads to their different shapes is not clear. Although it is known that slower cooling rates increases the amount of discontinuous precipitation, in the case of magnesium alloys other conditions and particularly alloy modifications that may prevent the formation of this detrimental morphology need be explored in more detail. 1.1

Discontinuous and Continuous Precipitation

In solid-solid phase transformations there are two reactions that lead to identical cellular morphologies, eutectoid transformation and discontinuous precipitation. Unlike the eutec-

152 toid transformation, discontinuous precipitation does not lead to decomposition of the matrix into two new phases. It merely changes the orientation and the composition of the matrix in which the precipitation of a second phase takes place. An important characteristic of this reaction is that the compositional change across the transformation front occurs discontinuously. The reaction mechanism of discontinuous precipitation has been reviewed by a number of workers and various suggestions have been made (18-21). Although not changing the general morphology, some exceptional examples of discontinuous precipitation have also been reported to exist (22-24). It should also be pointed out that apparently in all of the previous work the discontinuous precipitation has been investigated under isothermal reaction conditions rather than continuous cooling. Discontinuous precipitation in magnesium alloys has been studied in terms of composition profiles in cellular nodules (25-27). It was shown that the growth of discontinuous precipitation is affected by defects in matrix and the habit plane of the lamella can change from basal to pyramidal planes. It has also been recognized that regions of α phase behind the transformation front, though to a lesser degree compared to the untransformed matrix, remains supersaturated. In the case of discontinuous precipitation, it is known that the boundary diffusion predominates the growth of the precipitates and that supersaturation plays an important role. High angle boundaries are considered as sites of higher potential for the formation of the cellular nodules as is the case in magnesium alloys. However, the debate on a universally acceptable theory to allow the prediction of discontinuous precipitation is yet to be resolved. Continuous precipitation is, on the other hand, a more common phenomenon. Like its discontinuous counterpart, this reaction also gives rise to the formation of only one new phase. In systems where discontinuous precipitation is absent the term ‘continuous’ usually becomes redundant. This type of reaction does not involve a moving transformation front that is common to a colony of precipitates in the matrix grain. The precipitation merely takes place at energetically favorable intragranular sites such as dislocations, twin boundaries, and stacking faults generally by establishing an orientation relationship with the matrix. Each precipitate changes the composition of the matrix in their immediate vicinity, leading to a reduction in supersaturation, without changing the crystal orientation of the host lattice. However, the entire precipitation potential is generally not exhausted in continuous cooling conditions, or after short annealing times, and the precipitates that form first may not be the equilibrium type.

2

Experimental Procedure

The chemical composition of the alloy studied is given in Table 1. Preheated AZ91 scrap from a cold chamber die casting machine was heated to 720 °C in a steel crucible. A control sample was prepared by casting an ingot from the melt before nitrogen purging was started. The melt was then purged with nitrogen gas through a plunger for 30 minutes, and 60 minutes and ingots were cast representing both cases. All ingots were about 5 x 5 x 20 cm. in size. Table 1. Chemical composition of the alloy (wt%)

Mg

Al

Zn

Mn

Fe

Cu

Ni

Si

89.96

9.1

0.8

0.16

0.0061 0.0014 0.0010 0.005

Ca

Be

0.0068 0.0004

153 SEM examinations were carried out with a JEOL 840A microscope. TEM studies were conducted with a JEOL 2010 high-resolution microscope operated at 200 kV. Metallographic samples were cut from ingots by using a low speed diamond slit. Samples for light microscopy and SEM examinations were etched with 5% HNO3 solution. Thin foils for TEM were cut from a bulk specimen using spark erosion technique followed by mechanical polishing down to 100 μ, and dimpling with a micro grinder, and finally perforated via ion milling in cooled condition.

3

Experimental Results

Fig. 1 is a SEM micrograph showing the classical morphology in the sample not treated with nitrogen. Aluminum enriched regions towards the periphery of each grain were clearly noticeable. These regions are the result of coring as the equilibrium conditions were not achieved during the solidification of ingots. As shown in Fig. 1 the microstructure of the alloy was typical of AZ91 involving coarse β (Mg17Al12) particles along the grain boundaries as well as lamellar precipitation as a product of discontinuous cellular reaction.

Figure 1. SEM micrograph showing the microstructure of the control sample. Note the coarse β and the cellular precipitation along the grain boundaries and the contrast difference due to coring in grains.

Fig. 2 shows the microstructure of the ingot that was cast after purging the molten metal with nitrogen for 30min. As seen in the SEM micrographs the morphology of grain boundary coarse β particles is very different, displaying a jagged interface and copious amount of precipitation on them. Electron diffraction in TEM revealed that these are β particles of Mg17Al12 type (Fig. 3). TEM examination also identified the precipitates on β particles as αMg. No precipitation with cellular morphology was observed in this nitrogen treated sample. Fig. 4 is an SEM micrograph revealing the microstructure of the second ingot that was cast after purging the liquid metal with nitrogen for 1hr. The microstructure is clearly different from the other two samples. The coarse precipitates do not have a jagged interface nor the precipitates on them thus differing from those of the other nitrogen treated sample. On the other hand the cellular discontinuous precipitation product is also absent in the microstructure and therefore differing from the control sample that received no nitrogen treatment at all.

154

b Figure 2. (a) SEM micrograph showing the microstructure of the ingot cast after 30mins. of purging the molten metal with nitrogen. Note that the morphology of β has changed and that the cellular precipitation is absent. (b) a higher magnification SEM image of β particle.

a

b

(110)β (110)

Mg

c Figure 3. Bright Field (a) and Dark Field (b) TEM micrographs showing the coarse β; diffraction pattern (c) revealing the phases of β as Mg17Al12 and the precipitates as αMg.

155

Figure 4. SEM micrograph showing the microstructure of the ingot cast after 60 mins. of purging the molten metal with nitrogen. Note the absence of cellular precipitation and that the morphology of coarse β resembles that of the control sample.

4

Discussion

Knowing that the solubility of nitrogen in magnesium is practically nil one would perhaps not consider the addition of nitrogen as a potentially worthwhile trial towards creating a change in microstructure. The attempt on this ground may, at first sight, seem futile, or adventurous at best. However, this lack of solubility of nitrogen in magnesium coupled with the knowledge that the lamellar morphology is a product of an interface-controlled discontinuous reaction, are what led to this seemingly adventurous attempt. It was thought that if an unlikely small species such as nitrogen can be introduced into the alloy in the molten state, and if it can remain in solid at high temperatures in some minute quantities during solidification it would constantly be rejected to the crystal boundaries before it is totally diffuses out of the alloy. Assuming that such a scenario is valid one would then expect that introduction of nitrogen might alter at least the kinetics of the interface controlled reaction that forms the lamellar morphology and therefore modify it. The results of this study may be taken to support this proposition. Indeed, this study has shown that an ingot structure can be altered if the liquid metal is purged with nitrogen prior to casting. The period of purging seem to have different effects on the development of the microstructure. Shorter purging time seemed to promote a welldeveloped divorced eutectic leading to formation of coarse β particles with numerous αMg precipitations on them. This shorter nitrogen treatment seemed to have also changed the shape of the β particles leading to a serrated appearance along the edges. Longer purging appeared to have changed this appearance reverting back to the classical β morphology as seen in the control sample. Purging the molten metal prior to casting seems to change the microstructure considerably eliminating the cellular precipitation altogether. Admittedly, the explanation of the observed effects needs more detailed studies. Based on the current work the morphological differences in the appearances of β can only be related to the length of nitrogen purging operation.

156

6

References

[1] I.J. Polmear, Light Alloys, 3rd ed., Arnold, 1995, p. 196-247. [2] M.T. Murray, W.P. Sequeira, and R. D’Allesandro, SAE Technical Paper Series, 960420. [3] W.P. Sequeira, M.T. Murray, G.L. Dunlop, and D.H. John, Int. Conf. on Automotive Alloys, TMS, February, Orlando, Florida, 1997, p. 169-183. [4] H. Rosenson, Z. Koren, M. Bamberger, and E. Aghion, Proc. of the 9th International Metallurgy and Materials Congress, vol.2, Istanbul, Turkey, June-1997, p. 1537-1542. [5] A. Beste et al., Magnesium Alloys and Their Applications, eds. B.L. Mordike and K.U. Kainer, Proceedings volume sponsored by Volkswagen AG, Wolfsburg, Germany, April-1998, p. 253-258. [6] H. Sasaki, M. Adachi, and T. Sakamoto, Conf. Proc. IMA 53, Magnesium Material Advancing to the 21st Century, Ube City, Yamaguchi, Japan, June-1996, p. 86-92. [7] A. Mwembela and H.J. McQueen, ‘Hot workability of magnesium alloy AZ91’, Light Metals 1993, p. 523-532. [8] M. Dargusch et al., Proc. of 3rd Int. Magnesium Conf., edt. By G.W Lorimer, Inst. of Materials, Manchester, UK, April-1996, p. 153-166. [9] K. Pettersen, O. Lohne, and N. Ryum, Met. Trans. A, 1990, 21A, 221-230. [10] K. Pettersen and N. Ryum, Met. Trans. A, 989, 20A, 847-852. [11] E. Aghion and B. Bromfin, Magnesium Alloys and Their Applications, ed. B.L. Mordike and K.U. Kainer, Proceedings volume sponsored by Volkswagen AG, Wolfsburg, Germany, April-1998, p. 295-300. [12] J.B. Clark, Acta Metallurgica, 1968, 6, 141-152. [13] A.F. Crawley and K.S. Milliken, Acta Metallurgica, 1974, 22, 557-562. [14] A.F. Crawley and B. Lagowski, Met. Trans., 1974, 5, 949-951. [15] C.J. Bettles, ‘Magnesium Alloys and their Applications’, ed. B.L. Mordike and K.U. Kainer, Proceedings volume sponsored by Volkswagen AG, Wolfsburg, Germany, April-1998, p. 265-270. rd [16] P. Uzan, D. Eliezer, and E. Aghion, Proc. of 3 Int. Magnesium Conf., ed. G.W Lorimer, Inst. of Materials, Manchester, UK, April-1996, p. 43-50. [17] M.S. Dargush and G.L. Dunlop, Proc. of Materials 98, Institute of Materials Engineering, Australasia Ltd., ed. Michael Ferry, Wollongong, Australia, 1998, p. 579-584. [18] E. Hornbogen, Met. Trans., 1972, 3, 2717-2727. [19] W.Gust, Proc. Conf. on Phase Transformations, Inst. of Metals, 1979, II, p.27. [20] D.B. Williams and E.P. Butler, Int. Met. Rev., 1981, 3, 153-183. [21] G.R.Purdy, Proc.Int. Conf. on Solid-Solid Phase Transformations, ed. Aaronson, Laughlin, Sekerka, Wayman, AIME, 1982, p.521-530. [22] G.R. Purdy and N. Lange, 2nd Acta-Scripta Metallurgica Conf., Germany, 1983, p.214-221. [23] M. Kawase, H. Emoto and M. Kikuchi, Conf. Proc. on Phase Transformations, Inst. of Metals, Cambridge, 1987, p.254-258. [24] W.E. Voice and R.G. Faulkner, Conf. Proc. on Phase Transformations, Inst. of Metals, Cambridge, 1987, p.262-265.

157 [25] D. Duly, M.C. Cheynet, and Y. Brechet, Acta Metall., 1994, 42, 3843-3854. [26] D. Duly, M.C. Cheynet, and Y. Brechet, Acta Metall., 1994, 42, 3855-3863. [27] D.A. Porter and J.W. Edington, Proc. R. Soc. London A, 1977, 358, 335.

Recrystallization Behaviour of Four Magnesium Alloys L.W.F. Mackenzie+, F.J. Humphreys+, G.W. Lorimer+, K. Savage* and T. Wilks* +

Manchester Materials Science Centre, University of Manchester and UMIST, UK * Magnesium Elektron Ltd, Manchester, UK

1

Abstract

The thermomechanical behaviour of four magnesium alloys, AZ31, AS11, WE43 and ZC71, has been studied using channel die compression tests at temperatures between 558 and 673 K. The recrystallization behaviour of each alloy was observed between temperatures of 558 and 798 K, using optical microscopy and electron backscattered diffraction. AZ31, AS11 and ZC71 alloys underwent some dynamic recrystallization during the deformation at 558K. In AS11 and WE43 the nucleation of recrystallized grains occurred at both second phase particles and pre-existing grain boundaries; in AZ31 and ZC71 recrystallization was found to occur predominately at grain boundaries. The channel die compression results have been compared with the microstructures produced during extrusion at temperatures between 543 and 663 K.

2

Introduction

Some magnesium alloy extrusions are known to exhibit very different tensile and compressive strengths [1]. Extrusions with a fiber texture may have anisotropic mechanical properties if crystal orientation prevents twinning during tension parallel to the extrusion axis. This results in an increase in tensile strength relative to that in compression. It has been suggested that differences in texture in magnesium alloy extrusions can be attributed to recrystallization mechanisms and that particle stimulated nucleation of recrystallization (PSN) might be responsible for a reduction in texture [2]. The understanding of recrystallization behaviour is therefore of interest when considering mechanical properties. Study of recrystallization in magnesium alloys has been limited to a few systems such as Mg-Al-Zn and Mg-Zn-Zr [3, 4, 5]. It is important to develop an understanding of recrystallization mechanisms in magnesium alloys if the microstructure-property relationships are to be understood during thermomechanical processing. Four alloys were investigated in this work: AS11, AZ31, WE43 and ZC71.

3

Experimental

The alloys were received as sand castings of dimensions 200 x 200 x 25 mm. The sand cast AS11, AZ31 and ZC71 were homogenized at 708 K for 24 hours. In order to assist

159 the break-up of the second phase the ZC71 then underwent a 15% deformation by rolling at 673 K and a further 24 hour homogenization at 708 K. The WE43 was homogenized for 24 hours at 798 K and then 16 hours at 723 K. The alloys were deformed in a channel die to a true strain of 0.4 at a true strain rate of 10-4 s-1 and temperature of 558 K. 80 mm diameter billets of AS11, AZ31 and ZC71 were extruded into 19 diameter mm bars at 573, 555 and 543 K, respectively. WE43 was extruded into 38 mm diameter bar at 608 and 663 K. The tensile and compressive strengths of each alloy were measured parallel to the extrusion axis in the as-extruded state and after a recrystallization heat treatment. Full recrystallization was obtained when AS11, AZ31 and ZC71 were homogenized for 16 hours at 708 K and WE43 for 16 hours at 798 K.

4

Results and Discussion

4.1

AS11

Figure 1a shows the microstructure of AS11 deformed in a channel die at 558 K to a true strain of 0.4 at a true strain rate of 10-4 s-1. Twinning and nucleation of new grains occurred during deformation. Recrystallization is associated with original grain boundaries, second phase and eutectic particles. In Figure 1b recrystallization and twinning have occurred during extrusion at 573 K. New grains are associated with original grain boundaries, second phase and eutectic particles. Nucleation at small second phase particles

Nucleation at grain boundary

Twinning

Nucleation at particle

100 microns

Fig 1a. AS11 deformed in a channel die at 558 K to a -4 -1 strain of 0.4 at a strain rate of 10 s showing twinning and nucleation of new grains

50 microns Fig 1b. AS11 extruded at 573 K showing twinning and nucleation of new grains

The nucleation of new grains during deformation indicates that dynamic recrystallization has occurred in AS11 during both channel die deformation at 558 K and extrusion at 573 K. The association of new grains with second phase and eutectic particles indicates that dynamic PSN is probably a recrystallization mechanism in AS11. 4.2

AZ31

Figure 2a shows AZ31 deformed in a channel die at 558 K. New grains have nucleated at original grain boundaries in a “necklace effect” similar to that reported by Ion and Hum

160 phreys in a Mg-0.8Al alloy [6]. During extrusion at 555 K the AZ31 alloy fully recrystallised. Examination of the extrusion discard, Figure 2b, showed serrated grain boundaries and the nucleation of new grains similar to that reported by Galiyev, Kaibyshev and Sakai in a Mg-Zn-Zr alloy(4). New grains also nucleated at structural inhomogeneities such as shear bands. Fine grains nucleating at original grain b d i

100 microns Fig 2a. AZ31 deformed in a channel die at 558 K to -4 -1 a strain of 0.4 at a strain rate of 10 s showing nucleation of new grains

New grains nucleating at a serrated grain boundary

10 microns Fig 2b. AZ31 discard from material extruded at 555 K showing serrated grain boundaries beginning to nucleate new grains

The nucleation of new grains indicates that AZ31 underwent dynamic recrystallization during both channel die deformation at 558 K and extrusion at 555 K. Nucleation occurred predominantly at grain boundaries, although during extrusion new grains also nucleated at other structural inhomogeneities within the microstructure. 4.3

WE43

Figure 3a is of a WE43 sample deformed in a channel die at 558 K to a true strain of 0.4 at a true strain rate of 10-4 s-1 and then heated for 15 minutes at 673 K. The micrograph shows nucleation of new grains at original grain boundaries; many of the new grains are associated with grain boundary particles. New grains have also nucleated at particles within the matrix, similarly to that reported in an Al-Si alloy by Humphreys [7]. WE43 was extruded at 608 K and 663 K. When extruded at 663 K the material fully recrystallized. At 608 K some areas of the extrusion did not fully recrystallize, thereby allowing the identification of the origin of some new grains. In Figure 3b a new grain can be seen to have nucleated at a second phase particle. Figure 4 is an EBSD map of WE43 deformed in a channel die at 558 K and then partially recrystallized after a heat treatment of 15 minutes at 673 K. The thick and thin black lines represent high and low angle grain boundaries respectively. New grains have nucleated at original grain boundaries. Nucleation is also associated with second phase particles. A high concentration of low angle grain boundaries can be observed around some particles.

161

Nucleation at grain boundaries

Nucleation at particles

Nucleation at a particle

20 microns

10 microns

Fig 3a. WE43 deformed in a channel die at 558 K to -4 -1 a strain of 0.4 at a strain rate of 10 s and partially recrystallized at 673 K showing nucleation of new grains at grain boundaries and particles

Concentration of low angle grain boundaries at particle

Fig 3b. WE43 extruded at 608 K showing nucleation of a new grain at a particle

Nucleation associated with second phase particles

Nucleation at a grain boundary 50 microns Figure 4. EBSD map of WE43 deformed in a channel die at 558 K to a strain of 0.4 at a strain rate of 10-4 s-1 and partially recrystallized during a 15 minute heat treatment at 673 K

In both WE43 deformed in a channel die at 558 K and then partially recrystallized at 673 K and WE43 extruded at 608 K new grains were often associated with second phase particles. New grains also nucleated at grains boundaries and those new grains were often associated with grain boundary particles. This indicates that PSN is a mechanism of recrystallization in WE43. In Figure 4, EBSD has been used to highlight a concentration of low angle grain boundaries associated with a second phase particle; this could be associated with the formation of a particle deformation zone which is a precondition for PSN [7].

162 4.4

ZC71

Figures 5a and 5b show the microstructure of ZC71 deformed in a channel die at 558 K at a strain rate of 10-4 s-1. New grains have formed in a “necklace effect” at old grain boundaries. When extruded at 543 K the material was fully recrystallized with a very fine grain size; the origin of the recrystallized grains was not identified. Nucleation at grain boundary

20 microns 80 microns Figure 5a ZC71 deformed in a channel die at 558 K at strain rate of 10-4 s-1 showing nucleation of new grains at grain boundaries

Figure 5b ZC71 deformed in a channel die at 558 K at strain -4 -1 rate of 10 s showing nucleation of new grains at grain boundaries

The ZC71 alloy underwent dynamic recrystallization when deformed in a channel die at 558 K at a strain rate of 10-4 s-1. New grains were observed to nucleate predominantly at grain boundaries, Figure 5b. 4.5

Tensile and Compressive Testing

Each alloy was tensile and compression tested in the as-extruded form and then after heat treatment to produce a fully recrystallized microstructure. The values of 0.2 % proof stress in tension and compression are given in Table 1. AS11 and AZ31 alloys exhibited some anisotropy in tensile and compressive strength, which increased after recrystallization at 708 K. The WE43 and ZC71 alloys exhibited virtually no anisotropy in the as-extruded state. After heat treatment at 798 and 708 K, respectively, the WE43 again showed no anisotropy, whereas the ZC71 showed a large anisotropy. WE43 extruded at the higher temperature of 663 K exhibited no evidence of anisotropy in any condition. Table 1. Tensile and compressive strengths in the as-extruded and fully recrystallized conditions

Alloy AS11 as extruded AS11 recrystallized AZ31 as extruded AZ31 recrystallized WE43 as extruded WE43 recrystallized ZC71 as extruded ZC71 recrystallized

Extrusion Temp / K 573 543 608 555

Tensile Compressive Strength / MPa Strength / MPa 146 125 140 100 157 136 139 110 216 219 100 104 196 204 207 95

Tensile/Compressive Ratio 1.17 1.40 1.15 1.26 0.99 0.96 0.96 2.18

163 With fiber texture, crystal orientation allows twinning to take place during compression parallel to the extrusion axis but not in tension. This increases the tensile strength relative to the compressive strength, which may explain the anisotropy observed in AS11, AZ31 and ZC71. WE43 may have a random texture similar to that reported in previous research1.

5

Conclusions

The recrystallization behaviour of AS11, AZ31, WE43 and ZC71 has been studied during channel die deformation and hot extrusion. In AZ31 and ZC71 recrystallization occurred predominantly at grain boundaries, whereas in AS11 and WE43 new grains were observed nucleating both at grain boundaries and second phase particles; it appears that PSN occurs in these alloys. Tensile and compressive testing of extruded samples revealed anisotropy of strength in AS11 and AZ31 alloys, but not in WE43. The ZC71 alloy exhibited anisotropy in mechanical properties after full recrystallization but not in the as-extruded state.

6

References

[1] P. Prangnell, E. Ball, Scripta Met. And Mat., 1994, 31, 111-116. [2] F.J. Humphreys, M. Hatherley, Recrystallisation and Related Annealing Phenomena, 1st ed., Pergamon, Chapter 8. [3] M.R. Barnett in Magnesium Alloys 2003 (Eds.: Y Kojima , T. Aizawa, K. Higashi and S. Kamado), Trans Tech Publications, Switzerland, 2003. [4] A. Galiyev, R. Kaibyshev, T. Sakai in Magnesium Alloys 2003 (Eds.: Y Kojima, T. Aizawa, K. Higashi and S. Kamado), Trans Tech Publications, Switzerland, 2003. [5] X. Yang, H. Miura, T. Sakai in Magnesium Alloys 2003 (Eds.: Y Kojima, T. Aizawa, K. Higashi and S. Kamado), Trans Tech Publications, Switzerland, 2003. [6] S.E. Ion, F.J. Humphreys, S.H. White, Acta Metall., 1982, 30, 1909-1919. [7] F.J. Humphreys, Acta Met, 1977, 25, 1323-1344.

7

Acknowledgements

The authors would like to thank the EPSRC and Magnesium Elektron Ltd for their support.

Dynamic Recrystallization of Mg-3Al-1Zn T. Laser1, Ch. Hartig1, R. Bormann1, J. Bohlen2, D. Letzig2 1 2

Technische Universität Hamburg-Harburg, 21073 Hamburg, Germany GKSS Forschungszentrum Geesthacht, 21502 Geesthacht, Germany

1

Introduction

Due to their high specific strength, magnesium alloys offer a large potential in weight reduction to the automotive industry where less weight is leading to less emission and thus to more efficiency. The aim to fabricate structural parts out of Mg alloys drives the interest more and more to wrought alloys that offer a better ductility and strength compared to cast alloys. Magnesium wrought alloys can only compete against other light weight materials as long as the processing is inexpensive and well understood. Therefore, one step deformation of the feedstock that results in a fine recrystallized microstructure is a major goal. The poor cold formability of magnesium wrought alloys lets this goal only achieve at elevated temperatures accompanied with dynamic recrystallization (DRX). In several studies [1, 2, 3, 4] the dynamic recrystallization of Mg-3Al-1Zn (AZ31) during hot working has already been analyzed, but always starting from a fine grained microstructure (dm < 100 μm). In this work a coarse grained as-cast alloy was examined. The technological interest for a deformation process achieving a homogeneous microstructure with fine grain size was the main driving force.

2

Material and Experimental Procedure

The material employed in the present study was the wrought magnesium alloy AZ31. It was supplied in strand casting billets with 175 mm in diameter. The as-cast high purity material was cut in slices perpendicular to the cast axis and annealed for 20 hours at 400 °C for homogenization (air cool). Compression samples with an h/d-ratio of 1.5 were spark-eroded out of the slices with the compression axis lying parallel to the cast axis (Fig. 1). Uniaxial compression tests were performed at constant strain rates on a Zwick universal testing machine equipped with an electrical resistance furnace and a strain rate controlling unit. The tests were carried out in a temperature range between 200 °C and 400 °C in steps of 50 K. To ensure stable temperature conditions, the system with sample was held at testing temperature for approximately half an hour before deformation. The strain rate · -1 -2 -3 -4 -1 was varied in four steps: ε = 10 , 10 , 10 and 10 s . To reduce the friction between the sample and the compression plates, the top and bottom of the samples were lubricated with boron nitride. To conserve the microstructure that developed during forming, the samples were quenched in water immediately after unloading. The microstructure was studied by means of optical light microscopy (LM). With regard to uniaxial deformation only the center region of the compression samples was analyzed perpendicular to the load axis. The specimens were etched with an etching solution containing picric (effect on grain surface) and acetic (effect on grain boundary) acid.

165

Figure 1. Sample preparation out of strand cast billet (AZ31)

3

Figure 2. Initial strand cast microstructure (LM)

Results and Discussion

The initial as-cast microstructure after homogenization is shown in Fig. 2. An inhomogeneous grain size distribution with a mean grain size of 640 μm is observed. Precipitates of Mg17Al12 are homogeneously distributed in the matrix. 3.1

Flow stress behaviour

The true stress as function of the compression strain is shown in Fig. 3 for different strain rates and temperatures. The compression tests were carried out up to a maximum strain of ε = 1. In all flow curves the flow stress increases to a maximum and then decreases to a minimum for higher strains. This observed flow behavior is characteristic for hot working processes, where DRX leads to work softening [5]. The range of minimum flow stress is often termed in literature as a “steady state stress” [4, 6]. For very high strains (ε > 0.8) the flow stress may strongly increase especially at high flow stresses (i.e. low temperatures and high strain rates) due to friction. The maximum flow stress (σmax) moves to higher values of strain with increasing strain rate and decreasing temperature. At high temperatures and low strain rates, the maximum in the flow stress curve disappears and the flow behavior becomes nearly ideal plastic. The difference between the peak stress and the minimum stress is most significant for the mid temperature regime (250-350 °C).

Figure 3. Representative flow curves at different temperatures and strain rates

166 Neglecting the influence of friction, the minimum flow stress can be associated with a steady state flow stress σS [5,6]. Steady state flow commonly follows the Dorn equation: n

⎛˰ ⎞ ⎛ Q ⎞ ε& = A ⎜ V ⎟ exp ⎜ − m ⎟ G ⎝ ⎠ ⎝ RT ⎠

(1)

which has been used for a description of the measured steady state flow stresses as function of strain rate and temperature. Fig. 4 shows the strain rate as a function of the normalized1 stress for selected temperatures in a double logarithmic plot. The stress exponent · rises from n = 5.7 at 400 °C to n = 8.7 at 250 °C. At ε = 10-1 s-1 the resulting steady state flow stresses are lower than expected after a power law (Fig. 4). It has to be pointed out that creep laws may not be suitable for the description of processes where DRX is the controlling deformation mechanism superimposing the thermal activated slip of dislocations.

Figure 4. Power law behaviour in low strain rate regime, double logarithmic plot

Figure 5. Work hardening curves θ = ∂σ/∂ε, ε = 10 s , arrows indicating point of inflection ·

-3

-1

The work hardening behavior of the tests carried out at ε = 10-3 s-1 is shown in KocksMecking plots [7] in Fig. 5. After the theory of Poliak and Jones [8], a necessary condition for the onset of DRX is given by ∂/∂σ(-∂θ/∂σ) = 0, the point of inflection in the stage IV of the work hardening curve. After this condition DRX starts at stresses σ/σmax = 0.65 (T = 200 °C) up to σ/σmax = 0.9 (T = 350 °C) well below the peak stress, and at strains of ε ≈ 0.05, approximately 0.1 smaller than the corresponding peak strain (cf. Fig. 5). ·

3.2

Microstructural evolution

As already stated, all flow curves show a characteristic DRX behavior. This result can also be verified by examination of the microstructures. Figures 6a-6c show the microstructural evolu· tion for a specimen deformed at T = 300 °C and ε = 10-3 s-1 for different strains. All micrographs display the center of the specimen. From Fig. 6a it is obvious that the onset of DRX takes place at low strains before the peak flow stress is reached: fine globular DRX-grains have formed in the regions of grain boundaries. Some grains exhibit deformation twins. With increasing strain, several characteristic features develop in the microstructure (Figs. 6a-6c): (i) A flattening of the big original grains in load direction accompanied by further formation of 1

Shear modulus calculated with temperature dependence of Youngs modulus [9] for pure Magnesium.

167 deformation twins. (ii) At grain boundaries, necklaces of DRX grains are formed. Also at twin boundaries originating from deformation twinning, chains of DRX grains are visible at higher magnification (compare detail in Fig. 6b) (iii) At high strains, layers of DRX grains are clustering and consuming a major part of the microstructure. For strains ε ≥ 1.0, the necklace structure described above is always fully developed, i.e. all boundaries are decorated with chains of DRX grains. At low temperatures (Fig. 6d) deformation twinning increases drastically, fragmenting the globular grains of the initial microstructure. With an increase in temperature and a decrease in strain rate, the amount · of deformation twins decreases. At 400°C and ε = 10-4 s-1, no more twinning is observed.

a) T = 300 °C, ε& = 10-3 s-1, ε& = 0.15

b) T = 300 °C, ε& = 10-3 s-1, ε& = 0.3, detail showing onset of DRX at deformation twins

-3 -1 c) T = 300 °C, ε& = 10 s , ε = 1.0

-2 -1 d) T = 200 °C, ε& = 10 s , ε& = 1.0

Figure 6. Selected microstructures (LM), load axis from top to bottom

The total recrystallized volume fraction of each sample can be given qualitatively only because of overall inhomogeneous deformation conditions. A good quantitative approximation was found in limiting the region of analysis to the center region where uniaxial compression can be assumed. The results state that not more than 60% of recrystallized volume fraction can be obtained. This maximum of recrystallized volume fraction was · found for the mid temperature regime (T = 300 °C) and the lowest strain rate ε = 10-4 s-1. The mean size of the recrystallized grains situated in the DRX chains described above was determined. For a description of the grain size as a function of deformation conditions usually the steady state flow stress or the Zener-Hollomon parameter [6] are used. Fig. 7a shows the mean grain size dm as function of the normalized steady state flow stress σS for selected parameter sets. A mean grain size between 1 and 11 μm is observed under the cho-

168

a)

b)

Figure 7. Recrystallized grain size as function of a) normalized stress b) Zener-Hollomon parameter

sen deformation conditions. A clear flow stress dependence can be stated, the mean grain size decreases for higher stresses. For the deformation conditions illustrated by Figs. 6a-6c a mean grain size was observed being independent of the strain range (ε = 0.15 - 1.0). · The Zener-Hollomon parameter Z = ε exp(Q/RT) gives a reasonable guideline for deformation processes of technological relevance. For the determination of Z, a mean activation energy (Qm = 109 KJ/mol) was used, a value obtained from Arrhenius plots for constant stresses assuming power law behavior (Eq. 1). The mean grain size dm of recrystallized grains as a function of the Zener-Hollomon parameter is shown in Fig. 7b. The assumed slope changes for higher values of the Zener-Hollomon parameter, indicating different regimes of the evolution of DRX at high temperatures, low strain rates (low Z) and low temperatures, high strain rates (high Z). 3.3

Discussion

A qualitative similar behavior of the strain rate dependence of the flow stress was also observed for the magnesium alloy ZK60 [6] and for a fine grained magnesium alloy AZ31 [1] in compression tests. In detail, significant differences occur, especially concerning the strain hardening and the microstructural development (necklace structure). These differences may be ascribed to the coarse grained initial microstructure: The development of a necklace structure becomes much more obvious in coarse material and the resulting strain hardening is not only due to the competition between dislocation storage and dynamic recovery but also heavily influenced by the microstructural evolution [11]. An empirical relationship between strain rate, stress and temperature throughout the whole deformation range was found. Power law behavior was only observed for the low strain rate regime. At present, a complete description of the deformation behavior and microstructural evolution during DRX is lacking and only critical conditions for the occurrence of DRX can be given [8,10]. The results in the present study are in accordance with the “Poliak condition” (∂/∂σ(-∂θ/∂σ) = 0): The onset of DRX, as visible in LM, takes place after the “Poliak stress” but distinctly before the peak stress is reached. Similar as for ZK60 [6], at least two different regimes of the development of the grain size seem to be distinguishable, reflecting a change of nucleation mechanisms at low

169 temperatures and high stresses. A more detailed investigation of mechanisms for the material in the present study would require TEM observations, as done for ZK60 in [6]. The resulting necklace structure was also previously observed in AZ31 [2] and in ZK60 [6]. In [6] a theory for the formation of the necklace structure is proposed, where nonbasal slip of dislocations in the region of grain boundaries lead to accumulation of dislocations and hence to formation of high angle boundaries resulting in DRX grains. Tan and Tan [3] determined the optimum deformation conditions to receive a fine grained recrystallized structure in AZ31 to be T = 250°C at ε& = 10-4 s-1. In their study, a fine grained commercial material (12 μm) developed 85% of recrystallized volume fraction. For the coarse grained material analyzed in this study the optimum conditions (60% of recrystallized volume fraction) were found at the same strain rate but at a higher temperature: T = 300 °C. At 250 °C only 45% of recrystallized volume was observed. To receive high volume fractions of recrystallized areas, a two step deformation of the coarse grained initial microstructure might be advisable. In a first step a high fraction of deformation twins is generated at low temperature (200 °C) and low strain, whereas in a second step complete recrystallization at T = 300 °C may be achieved.

4

Conclusions

• The deformation of the studied coarse grained AZ31 alloy at elevated temperatures leads to the formation of DRX grains nucleating at grain and twin boundaries and forming a necklace structure throughout the whole deformation range. • A total recrystallized volume fraction of up to 60% can be obtained, generating DRX microstructures with mean grain sizes from 1 to 11 μm. • A rerystallized fine grained microstructure might be achieved in a two step deformation, producing high amount of deformation twins at low temperatures and high strain rates and a fully recrystallized microstructure at low strain rates in the mid temperature regime.

References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]

A. G. Beer, M. R. Barnett, Magnesium Technology 2002, 193-198 M. M. Myshlyaev, H. J. McQueen, Mater. Sci. Eng. 2002, A337, 121-133 J. C. Tan, M. J. Tan, Mater. Sci. Eng. 2003, A339, 81-89 A. Mwembela, E. B. Konopleva, H. J. McQueen, Scripta Mater. 1997, 37, 1789-1795 F. J. Humphreys, H. Matherly, Recrystallization and Related Annealing Phenomena, Pergamon Press, Oxford, 1995, 363-392 A. Galiyev, R. Kaibyshev, G. Gottstein, Acta mater. 2001, 49, 1199-1207 H. Mecking, in Work Hardening in Tension and Fatigue (Ed.: A. W. Thompson), TMS-AIME, New York, 1977, 67 E. Poliak, J. Jonas, Acta mater. 1996, 44, 127-136 F. Förster, W. Köster, Zeitschrift für Metallkd. 1937, 4, 116-123 G. Gottstein, E. Brünger, M. Frommert, M. Goerdeler, M. Zeng, Z. Metallkd. 2003, 94, 628-635 H. Mecking, G. Gottstein in Recrystallization of Metallic Materials (Ed.: F. Haessner), Dr. Riederer Verlag, Stuttgart, 1978, 195-222

Mechanical Properties and Formability of PM Mg-Al Based Alloys K. Matsuzaki, K. Hatsukano, K.Hanada, M.Takahashi and T. Shimizu National Institute of Advanced Industrial Science and Technology(AIST) AIST-EAST, 1-2 Namiki, TSUKUBA 305-3564, Japan

1

Introduction

Mg alloys are attractive for applications where weight reduction is required because of its density which is the lowest of all construction materials. Furthermore, the advantages of the Mg alloys are an unlimited resource and good process properties as well as good recycling. To this date, most Mg alloys are produced by casting processes such as diecasting or thixomoulding and the reaming part of those are produced by using plastic forming such as forging, press. The insufficient mechanical strength and poor formability of Mg alloys mainly restrict their application area. Also, the powder metallurgical (PM) method is a useful way to produce a high strength Mg alloy except for cost and safety of handling. The major benefits of the powder metallurgy are refinement of grain, minimization of segregations, homogeneous microstructure, extend of solubility, formation of nonequilibrium phase and so on. The high strength of above 600 MPa has been reported for the MgY-Zn alloy produced by PM method [1]. For a conventional Mg alloy such as AZ91, the tensile strength is improved up to above 300MPa by PM method [2, 3], which is three times higher than that of Ingot metallurgy AZ91. The fine microstructure introduced by PM process is main reason for the improvement of the strength and is also expected to improve the formability of Mg alloy. This paper intends to describes preparation of the AZ31 alloy by hot-extrusion of gasatomization powders, and clarify their mechanical properties and formability

2

Experimental Procedure

AZ31 alloy powders were produced by gas-atomization. The ingots of AZ31 alloys were melted in a graphite crucible by an induction furnace under an argon atmosphere and atomized at a melt temperature of 1143 K into powders using a mixture gas of Ar and 0.25%O2. The ejection pressure of the melt was held at constant value of 0.3 bar and the atomization pressure was 50 bar. The obtained powder has a spherical shape with some satellite, shown in Fig.1. The powder was sieved under particle sizes of below 75 μm, followed by hot-pressed into a perform with a diameter of 39 mm and a high of 40 mm in

171 air under a pressure of 50 tonn at 753 K in 10min, and then extruded into a bar at temperatures ranging from 548 to 623 K with an extrusion ratios of 1:10 and 1:2 0. Microstructure was examined by an optical microscope and a scanning electron microscope(SEM). Hardness was measured by a Vickers micro hardness tester with a 0.49 N load. Tensile properties (ultimate tensile strength, proof stress and elongation to fracture) were obtained by a tensile test with a -4 initial strain rate of 2.6 x 10 1/s at 50Ǵ㨙 room temperature. In order to examine Fig.1. SEM Photograph of gas atomized AZ31 powders the formability, a compression test was carried out to a deformation degree of 0.6 at room temperature and temperatures ranging from 373 to 673 K at an initial strain rate of 10-4 1/s. In the test, cylindrical specimens machined from the extruded bars were used. For comparison, mechanical properties and formability of a commercial wrought AZ31 alloy with a diameter of 25 mm produced by extrusion were also examined.

3

Results and Discussion

The hot-extruded sample consists of fine equiaxed grains and shows a homogeneous microstructure. The average grain size of the sample extruded at 673 K with an extrusion ratio of 1:20 is 10 μm and slightly increases with an increase in extrusion temperature and a decrease in extrusion ratio. The X-ray diffraction analysis reveals that the basal plain in hcp-Mg of hot-extruded AZ31 powders is slightly parallel to the direction of extrusion and no significant deference in the X-ray diffraction pattern is observed for the extrusion condition. The relative density of the extruded sample reaches 98%, indicating that a fully dense structure is obtained. Table 1 summarized the tensile mechanical properties at room temperature with a Vickers hardness of hot-extruded AZ31 powders. For comparison, the properties of a commercial wrought AZ31 alloy are listed. The sample produced by extrusion at 573 K with an extrusion ratio of 1:2 0 has a yield stress of 250 MPa and a ultimate tensile strength of 340 MPa, whose values are higher than those of the wrought AZ31 alloy. The strength decreases with decreasing extrusion temperature and extrusion ratio. The hardness is in the range of 75 to 78 HV. The increase in strength is mainly due to the refinement of grain size. The elongation increased with increasing extrusion ratio and extrusion temperature, and reaches 21% for the sample extruded at 623 K with an extrusion ratio of 1:2 0. This is maybe due to theconnection between the grains becoming stronger with increasing extrusion ratio and temperature, leading to the improvement of ductility. The SEM observation revealed that the fracture surface of the sample extruded at 623 K with

172 an extrusion ratio of 1:2 0 shows a ductile feature. On the other hand, the sample extruded at 448K with an extrusion ratio of 1:10 is composed of ductile and brittle regions. Table 1. Mechanical properties of AZ31alloys produced by PM

Yeld stress (MPa)

Ultimate tensile strength (MPa)

264

340

13

78

215

300

15

75

250

340

18

78

220 175

310 260

21 15

76 50

Extrusion condition Extruion Temp. 548 K Extrusion ratio 1:10 Extrusion Temp. 573 K Extruion ratio 1:10 Extrusion Temp.573 K Extruion ratio 1:20 Extrsion Temp. 623 K Extruion ratio 1:20 Wrought alloy

Elongation to Fracture (%)

Hardness (HV)

Therefore, an increase in extrusion ratio is expected to improve the strength and ductility. It is found that the hot-extruded AZ31 powder shows high strength at room temperature compared to the wrought AZ31 alloy. In order to evaluate formability, a compression test was carried out using a round specimen with a diameter of 8mm and a height of 10 mm at room temperature and elevated temperatures ranging from 373 to 573 K. Figure 2 shows the flow curves obtained from the compression test for the extruded AZ31 powders at 573 K with an extrusion ratio of 1:20. At room temperature, the sample shows an yielding and then the flow stress increases with deformation and finally fractures at 0.18. The yield stress and fracture stress is 220 MPa and 470 MPa, respectively. With an increase in deformation temperature, yield stress and flow stress decrease, and formability increases. At above 423 K, the deformation degree exceeds 0.6. Figure 3 shows the initial and compressed samples at various temperatures. The sample deformed at RT and 373 K shows a shear failure along a plane inclined at an angle of 45 °with respect to the compression axis. At above 423 K, the samples are deformed successfully to a deformation degree of 0.6 without defect. The samples extruded at deferent temperatures with an extrusion ratio of 1:20 show a similar deformability at above 423 K. This means that the sample with an extrusion ratio of 1:20 has good formability even at a temperature as low as 423 K. For the samples with an extrusion ratio of 1:10, however, deformed samples show a distinct deferent fracture feature as shown in Fig. 4, although the similar flow curves are observed. At room temperature, shear fracture along a plane inclined at about 45° to the compression axis and the cracking parallel to the compressive axis are observed.

Nominal compressive stress (MPa)

173

600

PM AZ31

500 400

Hor-extruded 573K, 1:20

373K

423K

R.T. 473K 523K

300 200

573K

100 0

0.1

0.2

0.3

0.4

0.5

0.6

Deformation degree Fig. 2. Flow curves at various temperatures for the AZ31 alloy produced by the hot extrusion of gas atomized powders

Initial

Compressed

R.T.

373K

423K

473K

573K 10mm

Fig. 3. External appearance of uncompressed and compressed AZ31 alloy prepared by extrusion at 573 K with an extrusion ratio of 1: 20

The cracking parallel to the compressive direction takes plces for the all samples with an extrusion ratio of 1:10 and occurs during the deformation even at 673 K, indicating the hot extruded AZ31 alloy powder with an extrusion ratio of 1:10 has a poor formability. SEM micrographs shown in Fig. 5 reveals that a region of shear fracture shows a smooth surface and in the region at the crakes parallel to the compressive direction the fracture takes places at interface between particle. The smooth shear fracture surface is observed for the sample with an extrusion ratio of 1:20 and the interparticle fracture surface is observed for the sample with an extrusion ratio of 1:10. These reasons are as follows. In the present study, the AZ31alloy powders were produced by gas atomization using an Ar gas containing 0.25% oxygen as atomization gas in order to passivate the powders. Thus, the powder seems to be covered with a rather thick oxide film. The extrusion with an extrusion ratio of 1:10 is insufficient to break the oxide films on the powder surface and a strong

174

Initial

Compressed

R.T

373K

473K

523K

573K 10mm

Fig. 4. External appearance of uncompressed and compressed AZ31 alloy prepared by extrusion at 573 K with an extrusion ratio of 1:10

connection between particles is not obtained. With a higher extrusion ratio, the powder is subjected to large shear deformation during the hot extrusion, leading to the formation of a new surface and thus a strong connection between particles. Therefore, it is said that the extrusion ratio of above 1:20 is needed to produce PM AZ31 alloys with good formability. One of the other methods to obtain the formability is to perform a successive PM operation from the powder production through consolidation under a clean atmosphere [4]. In the case of a wrought AZ31 alloy, a fracture occurs at a deformation degree 50Ǵ㨙 of 0.18 at room temperature and the deformation to a degree of 0.6 without cracks is achieved at temperatures above Fig. 5. Fracture surface of AZ31 alloy produced by hot523 K. This suggests that the formability extrusion of gas atomized powders at 573 K with an of PM AZ31 alloy is superior to that of extrusion ratio of 1:10 wrought AZ31alloy. The improvement of formability is also believed to be due to the refinement of grain Figure 6 shows the change in the hardness for the deformed PM and wrought AZ31 alloys as a function of deformation temperature. The hardness of 75 HV is obtained for the PM alloy after deformed at 423 K, at which the sample has an ability to deform to above 0.6 in deformation degree without defects. By increasing deformation temperature, hardness slightly decreases but retains 70 HV at 673 K. On the other hand, the wrought alloy after deformed at above 573 K shows a hardness of 40 HV. This means that the PM

175

4

After deformed 100

No cracks

Hardness (HV)

AZ31 alloy shows a high strength even after deformations. It allows us to produce a high performance Mg parts by PM method. Finally, it is said that the AZ31 alloy produced by PM processing shows a high strength at room temperature and good formability at elevated temperatures above 423 K. Moreover, the excellent formability at temperatures as low as 423 K is expected to make them useful in near net shape forming of high strength Mg alloys such as hot forging.

50

PMAZ31 573K,1:20

No cracks

Wrought AZ3㧝alloy 0

RT 373

423

473

523

573

Deformation temperature (K㧕 Fig. 6. Change in the hardness as a function of deformation temperature for the PM and Wrought AZ31 alloy

Conclusion

The gas atomized AZ31 powders was extruded into a bar. The obtained AZ31 alloy consists of fine grain with an average grain size of 10 μm, with a high tensile strength of 340 MPa and an elongation of 18% at room temperature. The strength slightly decreases but elongation increases to 21% with increasing extrusion temperature. The sample extruded at an extrusion ratio of 1:20 has good formability at temperatures above 423 K and deforms to a deformation degree of 0.6 without fracture. On the contrary, the decrease in extrusion ratio to 1:10 causes a cracks in the cylindrical surface during compression and reduces formability. The extrusion ratio plays an important role in formability. It is concluded that a high strength AZ31 alloy with good formability is obtained by hot-extrusion of gas-atomized powders at a high extrusion ratio.

References [1] Y. Kawamura, K. Hayashi, J. Koike, A. Inoueand T. Masumoto: Mater. Sci. Forum, 2000, 350-351. [2] G. Nussbaum, P. Sainfort, G. Regazzioni and H.Gjestland: Script Metall., 23 (1989) 1079-1084. [3] C. F. Chag, S. K. Das, D. Raybould, R. L. Bye and E. V.Limoncelli: Light Metal Age (1989) p. 12. [4] A. Kato, H. Horikiri, A. Inoue and T. Masumoto: Mate. Sci. & Eng., A179/180 (1994) 707-711.

Corrosion Properties of Conversion Plasma Coated Magnesium Alloys H. M. Nykyforchyn*, W. Dietzel**, M. D. Klapkiv*, C. Blawert** * Karpenko Physico-Mechanical Institute of the National Academy of Sciences of Ukraine, Lviv, Ukraine ** GKSS-Forschungszentrum Geesthacht GmbH, Geesthacht, Germany

Abstract If a voltage of several hundred volts is applied to a metal-electrolyte system, electric discharges of an intensity of nearly 107 V/m occur which result in a plasma developing in the discharge channels leading to oxide deposition on the metal surface. The oxide-ceramic coatings thus obtained on magnesium alloys are monophase in contrast to those on aluminium alloys, and the alloying elements contained in the magnesium essentially influence the microhardness and thickness of the coatings. In the tests reported here oxideceramic coatings having a thickness of 20 to 140 μm and a microhardness of 10 GPa were produced on the surfaces of three magnesium alloys. The corrosion resistance of these coatings was investigated in acidic (pH 3.1), near neutral (pH 7.1) and alkaline (pH 11.8 and 13.4) solutions using polarisation experiments. It was found that a change in pH from acidic to alkaline conditions shifted the corrosion potential from -1200…-1500 mV to 220…-280 mV. In general, the oxide-ceramic coatings created on magnesium alloys by an electrolyte plasma reaction decreased the corrosion current density by more than one order of magnitude in chloride containing solutions but practically had no beneficial effect in case of alkaline environment.

Introduction A major factor that limits the application of magnesium alloys is their comparatively low resistance against corrosion and wear [1]. One way to overcome this problem is the application of a suitable coating. Metal coatings however are of limited applicability since due to the extreme negative electrochemical potential of the magnesium they act as a cathode with respect to the magnesium substrate in the case of cavities or surface damages. The synthesis of oxides in a low-temperature electrolytic plasma allows it to cover surfaces of magnesium and its alloys with multifunctional oxide-ceramic coatings in the same manner as has previously been shown for aluminum alloys [2]. In an attempt to identify the underlying mechanisms of this synthesis and to optimise the technological process it is essential to know the composition and the parameters of the plasma and to further clarify the role of the oxides. In previous work the plasma luminescence at high potentials, i.e. 40-400 V, was investigated, and it was shown that in the sub-sparking domain this resulted from oxide luminescence whereas in the sparking domain where current

177 densities are less than 0.5 kA/m2 the luminescence is also induced by some electrolytes and by the base-metal constituents [3-5]. On the other hand, the results of these investigations can not directly be used to determine the most appropriate plasma producing parameters and/or to elucidate the role of the oxides. Moreover, the spectra of multicomponent electrolytes are what might be described as “contaminated”, which makes an exact identification of the source of the radiation impossible.

Experimental Procedure Three magnesium alloys were investigated having the compositions Mn 0.1-0.5, Al 4-6, Zn 1.3, Li 7-10, Cd 4 (designated “IMB2”), Mn 0.1-0.4, Al 0.5-1, Zn 2.3, Li 10-11.5 (“MA18”) and Zn 0.8, Y7.1-7.9, Cd 0.63, Zr 0.5 (“BMD10”), respectively. Oxide-ceramic coatings were produced by spark and microarc discharges using a set-up which is schematically shown in Figure 1 and which in principle consists of a current source and an electrolytic bath containing the specimen. Oxide-ceramic coatings were produced at cathodic to anodic current density relations of Ic/Ia = 1 using two different electrolytes: A KOH solution containing an addition of sodium silicate (solution “N1”) and a similar solution which in addition contained a chromium compound (solution “N2”). The thickness of the coatings thus obtained was varied from 20 μm to 140 μm by varying the process time. Figure 2 shows a magnesium specimen during the coating process surrounded by a plasma in the electrolyte bath at initial, intermediate, and final stage of the process, and a coated sample.

Figure 1. Principle of set-up (schematic, left) and time dependence of the voltage applied (right) for producing oxide-ceramic coatings: 1- specimen (working electrode), 2- electrolyte, 3- bath, 4- discharge channel.

The electron number density in the plasma, ne, was determined from the broadening of the spectral lines due to the linear Stark effect. The corresponding theory takes into account the broadening, Δλ, of the spectral lines caused by a slowly varying electric field and by fast electrons [6]. This theory ensures a high accuracy of the calculations and, due to the fact that the profiles of Balmer lines can be measured relatively easily, this method is convenient for the determination of the electron number density, ne. The complete method for calculating the electron number density and temperature is described elsewhere [7].

178 The corrosion properties of the thus obtained coated samples were investigated using a potentio-dynamic test technique and a chlorine-silver reference electrode. The corrosion environments consisted of 0.3% HCl solution with pH = 3.1, a 0.3% NaCl solution with pH = 7.1, and a KOH solution with pH-values of 11.8 and 13.4.

Figure 2. Magnesium specimen surrounded by a plasma in an electrolyte bath at initial, intermediate and final stage of the coating process (top to bottom, left), and coated sample (right).

Results and Discussion The spectrum obtained from an electrolyte plasma is a superposition of the continuous spectrum of electron radiation and line or band spectra of various atoms, molecules and radicals. The continuous spectrum appears as a result of free-free (bremsstrahlung radiation) and bound-free (collision-radiative recombination of electrons) transitions. The radiation energies of bound-free transitions are in the same order of magnitude as the radiation energies of spectral lines located near the edges of the series, and correspond to highly excited states of atoms close to the ionisation energy. These states are split according to the linear Stark effect. Micro-fluctuations of the field caused by the fields of neighbouring charged particles or by an external electric field broaden the corresponding levels and lines. The line emission spectrum (bound-bound transitions) characterises the electronic transitions in atoms and ions between certain energy levels and contains information on the physical parameters of individual atoms and ions and of the plasma as a whole. The kind of the corresponding elements and their degrees of ionisation can be determined from the wavelengths of the emission lines. The intensity of the lines and their half widths specify the temperature and the number density of the electrons. The experimental data shown in Figure 3 in terms of the relative radiation intensity, J, indicate that the entire emission spectrum consists of the continuous spectrum of the electron radiation and of individual lines of atoms and ions of magnesium, hydrogen, sodium,

179

J [relative units]

oxygen and the OH-radical. The parameters of the identified lines were used to determine the number density and temperature of the electrons [8]. The number density of the plasma electrons determined from the broadening of the Hα hydrogen lines was found to be ne ≈ 3.3×1022 m-3. In previous studies of oxide synthesis in an electrolyte plasma it had been assumed that the plasma temperature lies in the interval of 1000-2300 K [9]. However, the presence of partially fused-in spots in the coating leads to the assumption that the plasma temperature must be significantly higher. An analysis of the changes in the relative intensities of the spectral lines during the formation of the coatings with various thickness suggests that the plasma temperature is indeed much higher. The electron temperature determined from the ratio of intensities of the Mg I (383.8 nm) line and the Mg I (518.3 nm), Mg I (517.2 nm) Mg II (448.1 nm) lines was found to be 104 K. The maximum difference in the temperatures obtained for these lines taking the time dependence of the electron density into account was smaller than 400K.

λ [nm] Figure 3. Emission spectrum of an electrolytic plasma at a current density of 3 kA/m . 2

The analysis of the anodic and cathodic polarisation curves obtained in the potentiodynamic tests enables the determination of the potentials and corrosion currents of the 0 0 , icor ) and with the two coatings N1 alloys under investigation in the initial state ( Ecor 1 1 2 2 ( Ecor , icor ) and N2 ( Ecor , icor ) (E2cor, i2cor), respectively. These data are shown in Table 1. The corrosion potentials measured for the magnesium alloys are rather negative. In the NaCl solution they are shifted by 145 - 315 mV to negative values compared to the HCl solution. The corrosion rates of the uncoated magnesium alloys were higher in the 0.3% NaCl solution than in the HCl environment. Here, it should be pointed out that the concentration of Cl ions is about 6 times higher in the NaCl solution. This increase in Cl ion content locally destroys the oxide films promoting the so-called negative differentialeffect and hence increases the corrosion rate. Similar data have been reported for the corrosion of pure magnesium in chloride solutions [10]. The alloy IMB2 shows the highest corrosion resistance in both environments, despite of the fact that its potential is not the lowest. This may be due to the positive effect of Li, Cd, Zn as alloying elements. The

180 higher corrosion rate of the alloy BMD10 may be attributed to the presence of the cathodic phase of the cubic intermetallic alloy Mg24Y5. Table 1. Summary of the corrosion data obtained from potentio-dynamic tests. 0 Ecor mV -1377 -1365 -1209 -1556 -1518 -1522 -795 -964 -908 -449 -339 -499

Solution Alloy 0.3% HCl BMD10 IMB2 MA18 0.3% NaCl BMD10 IMB2 MA18 KOH, BMD10 pH11.8 IMB2 MA18 KOH, BMD10 pH13.4 IMB2 MA18

1 Ecor mV -1327 -1147 -1435 -1429 -1331 -1456 -829 -229 -587 -242 -280 -226

2 Ecor mV -1294 -1283 -1283 -1490 -1444 -1506 – – – – – –

0 icor , A/m 2 0.037 0.012 0.024 0.140 0.024 0.110 0.00611 0.00135 0.00344

0.0325 0.0239

1 icor , A/m 2 0.00263 0.00017 0.00388 0.00157 0.00125 0.01200 0.00408 0.00228 0.00528 0.0254 0.0143 0.0293

2 icor , A/m 2 0.00067 0.00120 0.00392 0.00240 0.00273 0.00101 – – – – – –

Note: The indices 0, 1 and 2 refer to alloys without coating and with coatings formed in the electrolytes “N1” and “N2”, respectively.

In general, the corrosion potential is becoming more noble with the oxide-ceramic coatings, and the corrosion rates of alloys with coatings decrease by 1-2 orders of magnitude. The extent to which the current density decreases is determined by the thickness of the coatings, their porosity, and further depends on the chemical composition of the alloy. Without coating and with the coating “N1”, the alloy IMB2 shows the highest corrosion resistance in both environments. With the coating “N2” the alloy BMD10 has the best resistance in the HCl solution, whereas the alloy MA18 has the best resistance in the NaCl solution. 0 Ecor, mV

a

-400

10

3

-1

6 2

5

1

-1200

icor, A/m

-800

4

-1600 2

4

6

8

10

12

14 pH

b

1

2

3

10

-2

10

-3

10

-4

5

2

6

4

4 2

6

8

10

12 pH

Figure 4. Dependence of potential Ecor (left) and current icor (right) on pH for the alloys MA18 (1, 4), IMB2 (2, 5), and BMD10 (3, 6) without coating (1, 2, 3) and with the coating N1 (4, 5, 6)

The effect of the coating thickness and the porosity on the corrosion current density is caused by anodic dissolution at the pore tips with associated with cathodic depolarisation of hydrogen at the side faces of the pores where the negative charges are accumulated.

181 The dependence of the potential and the corrosion current on the pH of the solution is not unequivocal (Figure 4): In the chloride environment the potential decreases and the corrosion current increases with increasing pH. On the other hand, in alkaline solutions the potential increases significantly and the corrosion current decreases (at pH 11.8) and which can be attributed to passivation processes. The positive effect of the coating is vanishing in this case. However, the corrosion current decreases again for pH 13.4 in spite of the potential also further increasing.

Conclusions The procedure outlined here allows to obtain oxide-ceramic coatings with a corrosion resistance in chloride containing environments which is by 1 to 2 orders of magnitude higher than that of uncoated material. The number density of the electrons was determined 22 -3 to be ne ≈ 10 m , and the electron temperature during the synthesis of oxides in electrolyte plasma of spark discharges is Te ≈ 104 K. An increase of the chlorine ion content in the environment decreased the corrosion resistance of the coated alloys. On the other hand, a passivation of the magnesium alloys in alkaline solutions brought the corrosion currents of the uncoated materials close to the levels which were observed for the same alloys with oxide-ceramic coatings.

References [1] H. Godard, W. Jepson, H. Bothwell, R. Kane, The Corrosion of Light Metals, Wiley, N.Y. 1967, p. 311. [2] M. D. Klapkiv, H. M. Nykyforchyn, V. M. Posuvailo, Surface and Coating Technology 1998, 100-101, 219. [3] S. Ikonopisov, A. Girginov, M. Maskova, Electrochim. Acta 1989, 5, 631. [4] K. Shimizu, S. Tajima, N Baba, Electrochim. Acta 1979, 24, 35. [5] A. Günterschultze, H. Betz;, Z. Phys. 1931, 4, 681. [6] H. R. Griem, A. C. Kolb, K. Y. Shen, Phys. Rev. 1959, 116, 4. [7] M. D.Klapkiv, H. M. Nykyforchyn, V. M. Posuvailo, Mater. Sci. 1994, 3, 333. [8] H. R. Griem; Plasma Spectroscopy, McGraw-Hill, New York 1963. [9] T. B. Van, S. D Bronn, G. P. Wirtz, Ceramic Bull. 1977, 56, 563. [10] R. Tunold, H. Holtan, M.-B. H. Berge et al., Corr. Sci. 1977, 17, 353.

The effect of variations in Melt Temperature on the Grain Structures in an AM60 die Casting Hans I. Laukli1, Otto Lohne1, Lars Arnberg1, Haavard Gjestland2, Stian Sannes2 1

2

Norwegian University of Science and Technology, Trondheim; Norsk Hydro ASA, Porsgrunn

Abstract The grain structures in cold chamber high pressure die castings (HPDC) of magnesium are complex and vary with position in the casting. A fraction of the metal solidifies in the shot sleeve before injection and the remaining melt rapidly freezes in the die cavity. The final grain structures can be described as a result of three different phenomena: a) Melt solidified in the die, b) fragments of dendrites formed at the sleeve wall and c) crystals formed close to the sleeve wall during pouring and subsequent settling. The initial melt temperature and the conditions prevailing in the shot sleeve affect the contribution of each mechanism to the final grain structure. In the present work the grain structures in thin walled AM60 die castings produced at three different melt temperatures and at various shot sleeve temperatures, have been investigated. The grain morphologies are observed to depend strongly on initial melt superheat and the three different solidification mechanisms can to different extent contribute to the final grain structure.

Introduction In the cold chamber high pressure die casting process (HPDC) a significant heat loss is experienced during filling of the melt into the shot sleeve. Occasionally all superheat has dissipated prior to the injection stage [1], thus initiating solidification. Shot sleeve solidification comprises the formation of a solid skin at the shot sleeve wall and equiaxed dendritic crystals in the melt. The skin can result in cold flakes in the casting [2]. A mixture of primary crystals, termed externally solidified grains [3] or floating crystals [4], and residual melt is injected into the cavity. The relatively cold die walls and thin walled geometry ensures rapid freezing which generates a bimodal grain size distribution [4-7]. The general understanding of floating crystals have been limited to being recognized as externally solidified crystals. The mechanisms by which the crystals are formed in HPDC have been questioned [6]. A review of the different mechanisms of nucleation of equiaxed crystals in aluminium alloy castings has been reported [8], and analogous mechanisms have been suggested to pertain to the solidification of magnesium alloys [9]. Generally, the solidification of alloys initiates with crystal nucleation on suitable substrates [10] at a sufficient undercooling of the melt. These crystals can then grow into dendrites with different order branches depending on the dendrite tip radius, the local thermal gradient, solute liquid diffusion and the constitutional conditions present. When

183 the crystals start to interact, i.e. the thermosolutal diffusion fields encapsulating the crystals impinge on each other, peripheral growth is substituted by crystal coarsening [11]. The mechanisms of nucleation of equiaxed crystals in the absence of grainrefiner additives are briefly as follows: Equiaxed dendritic crystals can form during pouring of metal into a permanent mould, as a result of nucleation and growth in the thermally undercooled region adjacent to the wall [12]. These crystals grow with a necked shape due to the constraints set up by the wall restricting solute diffusion at the root. The crystals are easily separated from the wall by convection. A similar mechanism is the nucleation and growth of equiaxed crystals, termed free chill crystals, in the thermally undercooled zone adjacent to the mould walls when pouring into a cold permanent mould [13]. During growth of solid at the wall, solute is rejected at the crystal interface creating a solute enriched zone. This results in a region with lower liquidus temperature than the bulk of the liquid. Simultaneously, heat is extracted from the growing solid and a thermal gradient is established adjacent to the growing interface. A constitutionally undercooled zone is therefore generated in which suitable substrates can nucleate crystals [14]. When a layer of material solidifies at a mould wall, dendrites are rapidly formed growing into the liquid opposite and parallel to the direction of heat flow. The dendrite trunks develops branches in which segregated melt is interdendritically entrapped. These dendrites can be fragmented through remelting due to constitutional gradients, thermal conduction, solute solid diffusion or ripening phenomena [15]. The dendrite fragments then float or sink depending on the alloy constitution and convective conditions present. Another proposed mechanism is that due to heat radiation at the melt surface a thermally undercooled zone is generated adjacent to the melt surface in which dendrites can form. These crystals can be detached by breaking off [16]. A critical criterion for the survival of crystals is the magnitude of superheat in the melt. The present paper presents an investigation of the influence of melt temperature on the grain structures in thin walled AM60 box-shaped castings [17] produced in a Bühler 420ton cold chamber HPDC machine.

Experimental Approximately 15 boxes were cast to obtain a quasi-steady state temperature in the shot sleeve. These initial castings were thus produced with constant melt temperatures at the initiation of filling, at qualitatively increasing shot sleeve temperatures. At steady state a minimum of 30 castings were produced for each melt temperature with process parameters as listed in Table 1 (details in [6]). The steady state shot sleeve temperature has been estimated to be 250 °C directly below the pouring hole and 300 °C at the end [6]. Table 1. Process parameters used for producing AM60B die castings. (Tliquidus of the AM60 alloy is 619 °C o

o

-1

Tm [ C]

Tdie [ C]

Fill fraction

v2 [ms ]

P3 [bar]

640, 680, 710

200

~0.5

4.5

600

[18]

).

Samples were prepared for optical microscopy [19] taken from two positions in front of one of the gates, A) close to the gate and B) farther away, and C) in an area midway between the gates in the bottom of the box [4, 6]. Area fraction measurements of the floating crystals were performed with an equivalent circle diameter (ECD) threshold of 10 μm, derived from image analysis [20].

184

Results Effect of melt and shot sleeve temperature on the grain size distribution The micrographs in Figure 1 are obtained from a casting produced with an initial melt temperature of 640 °C. The microstructure is generally coarser closer to the gate (position A), with the primary crystals distributed over the cross section. In position B the coarse crystals are principally centered with fewer and finer floating crystals distributed from the center to the surface. A more random distribution and fewer coarse crystals are present in position C.

Figure 1. Optical micrographs of etched cross sections (2.5mm thickness) of a casting produced with low superheat, Tm = 640 °C. Micrograph A: flow vertical. B: flow horizontal. C: flow horizontal.

The measurements of area fraction as function of thickness for the three melt temperatures are displayed in Figure 2. Consistent with the micrographs in Figure 1, the area fraction of floating crystals is decreasing with distance from the gate. In position A (640 °C), the floating crystals covers 35% of the area in 50% of the cross section thickness. At 680 °C this drops to 30% in 30% of the thickness. At 710 °C the area fraction of floating crystals is approximately 10% in 10% of the section thickness. A maximum in area fraction is present in the center of position B for all temperatures. At 640 °C the area fraction is approximately 40%, at 680 °C it is 35% and at 710 °C 13%. A slight increase in area fraction is observed closer to the surface at 640oC. In position C a maximum in area fraction is observed at 640 °C, i.e. 25%. At 680 °C and 710 °C the area fraction in position C is approximately 2-3%. A C

0 Tm=640oC

10

20

30

40 50 60 [%thickness]

70

80

90 100

45 40 35 30 25 20 15 10 5 0

A B C

0 Tm=680oC

10

20

30

40 50 60 70 [%thickness]

80

90 100

45 40 35 30 25 20 15 10 5 0

A

Areafraction [%]

Areafraction [%]

B

Areafraction [%]

45 40 35 30 25 20 15 10 5 0

B C

0 Tm=710oC

10

20

30

40 50 60 [%thickness]

70

80

90 100

Figure 2. Area fraction of floating crystals in A, B and C for Tm = 640 °C (left), 680 °C (center) and 710 °C (right).

In Figure 3 micrographs from position B from a) a pre-steady state casting and b) a steady state casting are displayed. A coarser microstructure is present in Figure 3 a). a)

b)

Figure 3. Micrographs from position B. Tm = 710 °C. a) Lower Tshot sleeve and b) at a quasi-steady state Tshot sleeve.

185 Effect of superheat on the crystal morphology The castings produced at a low superheat (640 °C) revealed two distinct crystal morphologies. The floating crystals are observed to be coarse dendritic or rosette like, as observed in Figure 4 a), or globular in shape, Figure 4 b). The latter dominates in position C. At higher melt temperatures (680 and 710 °C) the crystal morphology is generally elongated dendrite trunks or arms (Figure 4 c) or branched dendrites (Figure 4 d). The branched crystals dominate in the center of the cross sections and closer to the gate.

Figure 4. Floating crystal morphologies; a) coarse dendritic/rosette like, b) globular (Tm = 640 °C), c) elongated trunks/arms, d) branched dendritic (Tm = 680 and 710 °C).

Discussion Effect of melt and shot sleeve temperature on the grain size distribution In magnesium HPDC up to 20% of the metal can be solidified in the shot sleeve prior to injection as reported in [5], and a decreasing fraction of floating crystals from the gate has been observed [21]. This is analogous to the observations in Figure 1 and in Figure 2 and has been proposed to result from the non-isothermal conditions in the shot sleeve. Supported by work on plunger movement and wave behavior in the shot sleeve, the larger fraction of coarse crystals observed close to the gate in the casting has been proposed to originate from closer to the plunger. The lower fraction of crystals observed farther from the gate has been deduced to originate from closer to the runner end [4, 6, 18]. Presolidification is most likely susceptible to shot sleeve volume and fill ratios. Measurements of the area fraction of floating crystals (with an ECD threshold of 6.7μm) close to 0.5 has been reported [5] using a small diameter shot sleeve with a fill fraction of 0.35. It can be deduced that with a low melt temperature a larger undercooling can be rapidly established and a larger number of crystals are nucleated (640 °C, Figure 2). Nucleation, growth and remelting of crystals depend on the degree of melt undercooling, and at higher melt temperatures the undercooling adjacent to the walls is reduced and fewer crystals are formed. Crystals that are nucleated later remelt due to the superheated liquid and hence fewer floating crystals are observed in the casting (Figure 2, 680 °C and 710 °C). The floating crystals are occasionally centered in the cross section thickness, e.g. position B in Figure 2, which can be explained by a solid skin is rapidly formed at the die walls during filling. The skin alters the flow behavior locally by reducing the effective cross section thickness. The melt, or mush, arriving later is forced towards the center. Thus, a larger area fraction of floating crystals is observed in the center [6, 18]. Occasionally, the skin may not be sufficiently thick to fully accommodate for the centered crystals and crystal migration can be an additional mechanism acting as proposed in [22]. Presently it is not clear if the observed plug flow is solely due to any one of these mechanisms, and work is in progress to increase the understanding of semisolid flow.

186 It has been proposed that velocity and superheat is critical for the distribution of floating crystals in the cross section thickness [6]. The microstructure in position B (Figure 1) deviate from that reported and can partially be attributed to the different filling phase velocities, i.e. 4.5 ms-1 in contrast to 6.5 ms-1 [6]. A lower velocity can result in a greater fraction of metal freezing at the die walls during filling. Hence, the crystals distributed between the surface and the center are locked in this position. This can not be observed at higher melt temperatures as fewer floating crystals survive in the shot sleeve, and the crystals that first enter the cavity are less susceptible to be deposited closer to the surface as less metal freeze at the die walls. The observed effect of decreased fraction of floating crystals with increased shot sleeve temperature can be deduced to result from the reduced number of crystals formed and the increased likelihood of remelting with superheated melt (Figure 3). To test the hypothesis that the thermal conditions in the shot sleeve affects the formation of floating crystals, a few castings were produced with Figure 5. Micrograph from position A coating in the interior of the shot sleeve. The coating was in a casting produced with coated shot sleeve. The coating procedure inter- expected to exaggerate the effects observed in Figure 3. rupted steady state and hence gener- As depicted in Figure 5, only a few floating crystals ated a slightly lower shot sleeve tem- are present in position A (close to the gate) despite that perature. Initial T = 680 °C. a longer cycle time was inevitably generated during coating. m

Effect of superheat on the crystal morphology The diversity of dendrite morphologies observed is proposed to result from the variable cooling conditions and different nucleation mechanisms that prevail in the shot sleeve. At lower melt temperatures a less inclined temperature gradient and a larger thermal undercooling is established adjacent to the shot sleeve wall in which crystals can nucleate. During filling the crystals can be formed at the wall and eventually separate according to the theory outlined in [12], or resemble the nucleation of free chill crystals in the thermally undercooled region adjacent to the shot sleeve wall [13]. The “wall crystals”, with a nearglobular shape [12], are transported into the bulk of the liquid during filling by the flow of melt. The crystals survive as the melt is sufficiently cold and branching is restricted due to lack of undercooling. At a later stage a large number of floating crystals has formed and grown to the point of impingement followed by dendrite coarsening, which results in rosette like dendrites (Figure 4 a). The first wall crystals that were formed possess a globular morphology (Figure 4 b) and are distributed closer to the runner end of the shot sleeve and therefore end up far from the gate in the casting [4, 6, 18], e.g. in position C (Figure 1, right micrograph). As the globular crystals are associated with a lower melt temperature and dominates at lower solid fractions at low superheat, e.g. position C, they are most likely not formed as a result of the high shear rates encountered in HPDC, as proposed in [5]. To obtain a globular microstructure shear is commonly applied for order of magnitudes longer times than the crystals are exposed to in HPDC [23].

187 At higher melt temperatures the temperature gradient adjacent to the shot sleeve wall during filling is steeper, and wall crystals are not expected to form as easily as the thermally undercooled region is reduced. Some crystals can form, but these are disposed to shrink and remelt as they are carried into the bulk of the superheated liquid. At a later stage crystals nucleate on the wall and grow as dendrites into the melt. Dendrites can be nucleated in the constitutionally undercooled zone at the growth front [14] or be fragmented due to dendrite arm remelting at the roots [15]. The latter mechanism is more probable as the morphology of many of the crystals resembles dendrite fragments (Figure 4 c and d). A cascade effect can be triggered as more fragmentation is initiated. The dendrite finds itself in superheated liquid, and arms are easily detached. In-situ solidification experiments, by means of X-ray imaging, have manifested the ease by which dendrite arms are fragmented during columnar dendritic growth in aluminium alloys, and convection is a major contributor in the fragmentation process [24]. Forced convection generated by the plunger movement can enhance fragmentation through dendrite arm remelting. Forced and free convection is set up by the melt flow during and after pouring and thermosolutal gradients, respectively. The crystals remain branched dendritic till injection is initiated, as an insufficient number of crystals are present to impinge on each other and generate coarsening. It has been reported that dendrite fragments results from the high shear rates encountered in the gate [5]. If this mechanism pertains, a high fraction of branched dendritic crystals would be expected in the runner system. However, a large fraction of dendrite fragments is present on either side of the gate, as observed in the micrograph in Figure 6 obtained from a position in front of the gate. Metallic dendrites are not expected to break mechanically in a casting operation, but can be plastically deformed [25]. Figure 6. Microstructure of gate-section with dendrite fragments present.

The solidification in HPDC is highly non-idealised. Large differences in grain sizes, distributions and crystal morphology can be observed, depending on position in the casting, process conditions, thermal and constitutional variations, flow behaviour, volume fraction of pre-solidified metal and alloy type. An additional aspect can be the growth of floating crystals in the cavity, which is not within the scope of the present work. However, it can be argued that a higher melt temperature would generate a coarser microstructure due to a higher fraction liquid. Analogous, a low melt temperature and larger fraction of floating crystals would generate less melt to freeze, and coarsening would not occur. However, the floating crystals are coarser in the latter case. A higher melt temperature can improve the efficiency of the intensification pressure, and pressure effects (ClaussiusClapeyron) could be introduced generating a finer grain structure.

Conclusions The initial melt temperature and the thermal conditions in the shot sleeve affect the distribution of floating crystals:

188 1. A larger area fraction of floating crystals is obtained with a low superheat and colder shot sleeve, which is attributed to the large number of floating crystals formed and the increased survival rate. 2. A lower area fraction floating crystals is obtained with a larger melt superheat or altered heat transfer conditions in the shot sleeve. This is proposed to result from fewer floating crystals are formed and fewer survive. The floating crystal morphology is affected by the melt superheat: 1. With a low melt temperature wall mechanisms are proposed to dominate. This results in floating crystals with a coarse dendritic (rosette like) or globular morphology. 2. Branched and elongated dendrites are obtained with higher melt temperatures. This is attributed to extensive dendrite fragmentation in the shot sleeve.

Acknowledgements The present work was funded by the project NorLight Shaped Castings and partners: the Norwegian Research Council, Alcoa Automotive Castings, Scandinavian Casting Center ANS, Elkem Aluminium ANS, Fundamus AS, Hydro Aluminium Metal Products, Norsk Hydro ASA, NIMR, NTNU and SINTEF (project responsible). The authors thank the partners for financial support.

References [1] Z.W. Chen, M.Z. Jahedi, Int. J. C. Met. R., 1998, 11, 129-138 [2] D. Rodrigo, M. Murray, H. Mao, J. Brevick, C. Mobley, R. Esdaile, NADCA Trans., 1999, pp. 219-225. [3] D. StJohn, A.K. Dahle, T. Abbott, M.D. Nave, M. Qian, Proc. Magnesium Techn., 2003, pp. 95-100. [4] H.I. Laukli, O. Lohne, H. Gjestland, S. Sannes, L. Arnberg, NADCA Trans., 2002, pp. T02-035. [5] W. Sequeira, PhD thesis, The University of Queensland (Brisbane), 2000. [6] H.I. Laukli, O. Lohne, S. Sannes, H. Gjestland, L. Arnberg, Int. J. C. Met. R., 2003. Submitted. [7] A. Bowles, J.R. Griffiths, C.J. Davidson, Proc. Magnesium Techn., 2001, pp. 161168. [8] J. Hutt, D. StJohn, Int. J. C. Met. R., 1998, 11, 13-22 [9] A.K. Dahle, Y.C. Lee, M.D. Nave, P.L. Schaffer, D.H. StJohn, J. Light Met., 2001, 1, 61-72 [10] L. Bäckerud, M. Johnsson, Proc. Light Met., 1996, pp. 679-685. [11] W. Kurz, D.J. Fisher, Fundamentals of Solidification, 1998, pp. 85-89. [12] A. Ohno, Solidification - The sep. theory and its practical applic., 1987, pp. 51-60, 78. [13] B. Chalmers, J. Aust. Inst. of Metals, 1963, 8, 255-263 [14] W.C. Winegard, B. Chalmers, Trans. ASM, 1954, 46, 1214-1224

189 [15] K.A. Jackson, J.D. Hunt, D.R. Uhlmann, T.P. Seward, Trans. AIME, 1966, 236, 149158 [16] R.T. Southin, Trans. AIME, 1966, 239, 220-225 [17] S. Sannes, H. Gjestland, H. Westengen, H.I. Laukli, Proc. Mg. Alloys & Appl., 2003. [18] S. Sannes, H. Gjestland, H. Westengen, H.I. Laukli, O. Lohne, SAE Trans., 2003, pp. 03M-192. [19] H.I. Laukli, O. Lohne, L. Arnberg, Poster in Proc. Mg. Alloys & Appl., 2003. [20] Adobe Photoshop, 7.0, Adobe, 2002 [21] P.D.D. Rodrigo, V. Ahuja, Proc. Int. Conf. Mag. Sci. Techn., 2000, pp. 97-104. [22] A.K. Dahle, S. Sannes, D.H. St. John, H. Westengen, J. Light Met., 2001, 1, 99-103 [23] Y. Ito, M.C. Flemings, J.A. Cornie, Rheol. & microstr. of AlSi6.5 in Nature and prop. of SSM, 1991. [24] R.H. Mathiesen, L. Arnberg, K. Ramsøskar, T. Weitkamp, C. Rau, A. Snigirev, Met. Mat. Trans. B, 2002, 33B, 613-623 [25] A. Hellawell, S. Liu, S.Z. Lu, JOM, 1997, 18-20

Microstructural Characterisation of AZ91 Magnesium Alloy Andrzej Kiełbus*, Maria Sozańska*, Lubomirc Ciżek** * Silesian University of Technology, Katowice, Poland **VSB-TU Ostrava, Czech Republic

Abstract Casting magnesium alloys make up about 85÷90% of all products in Europe fabricated with the application of Mg. The most often used alloys are Mg-Al-Zn-Mn-based ones. Magnesium alloys are widely used in the automotive industry, particularly for elements of the car interior, car body, chassis and driving gears. The basic alloy, which individual motorcar elements are made of, is AZ91. The paper presents microstructural charakterisation of AZ91 (Mg-9Al.-0,6Zn) alloy in its initial state (after casting). Melting was performed in a 250 kg induction furnace by applying modification with SPEFINAL T 200 and refining with Emgesal Flux 200. The casting temperature was 740 °C. The research methodology covered microscopic investigation as well as LM, SEM and an X-ray microanalysis.

Introduction General characterisation of magnesium alloys Magnesium and its alloys are more and more often applied in different spheres of life, including the aircraft and motor vehicle as well as metallurgical, chemical and electricalchemical industries. This is connected, among other things, with the advantageous dependence between tensile strength of 160 to 365 MPa, modulus of elasticity (45 GPa) and density (1.74 g/cm3). Magnesium alloys have a high value of tensile strength/density compared to other alloys. Moreover, magnesium shows relatively good electric and thermal conductivity as well as high damping capacity [1, 2]. The basic magnesium alloys include ones which contain manganese, aluminium, zinc, zirconium and rare-earth elements which allow obtaining suitable properties. Manganese does not cause any increase of tensile strength, however, it does slightly increase the yield point. It also brings about an increase of resistance to the action of sea water. The quantity of manganese in magnesium alloys is limited by its relatively low solubility in magnesium. Manganese content in alloys with an Al addition does not exceed 0.3% and 1.5% in alloys without Al addition. Aluminium enhances both tensile strength and hardness, and improves casting properties of an alloy. The best ratio of mechanical to plastic properties is obtained with a 6% Al content. An addition of zinc in combination with Al aims at improving tensile strength at a room temperature, however 1% of Zn with a 7÷10% Al content in an alloy enhances hot cracking. Zirconium is added to alloys which contain

191 zinc, rare-earth elements, thorium and their combinations, for the purpose of structure refinement. It should not be used in alloys containing aluminium and manganese, since it forms stable compounds with them which are removed from the solid solution. Rare-earth elements are added to manganese alloys as a mischmetal or didymium. Mischmetal contains cerium, lanthan and neodymium, whereas didymium is a mixture of neodymium and praseodymium. An addition of rare-earth elements enhances magnesium alloys’ strength at a room temperature and what is more, it reduces porosity of casts [3]. In general, magnesium alloys can be divided into two groups [1, 3]: I – those containing 2-10% Al with minute quantities of Zn and Mn. They cost of production is low and they are characterised by good corrosion resistance as well as a considerable decrease of mechanical properties with an increase of temperature. The following alloys are rated among this group: • Mg-Al-Mn alloys; • Mg-Al-Zn-Mn alloys; II – those containing a wide range of elements (mostly Mn, Zn, Th, Ag and Si instead of Al), but always with an adequate content of Zr which has a considerable influence on the increase of mechanical properties. These alloys can work at elevated temperatures, however, the price of alloying additions along with a special production technology considerably raises the cost of their fabrication. Within this group of alloys, those most often applied are: • Mg-Zr alloys; • Mg-Zn-Zr alloys; • Mg-rare-earth elements-Zr alloys; • Mg-Ag- rare-earth elements-Zr alloys; • Mg-Y- rare-earth elements -Zr alloys; Casting magnesium alloys make up about 85÷90% of all products fabricated in Europe with the application of Mg. Those most frequently applied are AZ91 and AZ31 alloys which contain 9 and 3% Al respectively, and ca. 1% Zn; next are AM50 and AM60 alloys with 5 and 6% Al contents and with a Mn addition. These alloys are characterised by good casting properties, in particular during high-pressure casting, as well as good mechanical properties. They are used for casting such structural components of cars as: wheel rims (Fig. 1a), instrument panels, steering wheels (Fig. 1b) and seat frames (Fig. 1c), as well as components of equipment in other fields, e.g. telephones (Fig. 2a), camcorders (Fig. 2b) and gardening tools (Fig. 2c). a)

b)

c)

Fig. 1. Car components made of magnesium alloys: a) Ford wheel rim of AZ91D alloy; b) steering wheel of AM50A alloy; c) seat frame of AM60B alloy.

192 The maximum solubility of aluminium in magnesium at an eutectic temperature (437 °C) is 14%, whereas an eutectic mixture (α + Mg17Al12 intermetallic phase) occurs at ca. 33% Al content (Fig. 1). The content of aluminium in all industrial alloys of Mg with Al is not higher that the boundary solubility of Al in Mg. The equilibrium structure of these alloys is characterised by 100% presence of a solid solution, whereas the unbalanced structure, additionally metastable in casting alloys, shows the presence of an eutectic already at a 2% Al content [4]. a)

b)

c)

Fig. 2. Components used in different sectors of economy, made of magnesium alloys: a) fragment of a cell phone made of AZ91D alloy; b) camcorder casing; c) garden shears made of AZ91D alloy. 700

10

0

20

Magnesium, wt% 40 50 60

30

70

80

90

660°C

650°C

L

600

500 (Al)

450°C 18.6

458°C

452°C

λ

410°C

400

Mg17Al12

437°C 62.3

88.5

69

(Mg)

ε Al3Mg2

Temperature, °C

100

300

200

100

0

Al

10

20

30

40 60 50 Magnesium, at%

70

80

90

100

Mg

Fig. 3. Mg-Al binary phase diagram [3]

Experimental procedure Material for research The object of the research was the AZ91D casting alloy intended for wheel rims, cast into a sand mould. The casting practice included melting in a 250 kg induction furnace with the application of modification with SPEFINAL T 200 as well as refining with Emgesal Flux 200. The casting temperature was 740 °C. The chemical composition of the analysed alloy is presented in Table 1.

193 Table 1. Chemical composition of AZ91D alloy in wt.-%

Mg

Al

Zn

Mn

Si

Cu

Fe

Be

Remainder

9,15

0,6

0,24

0,03

0,01

0,01

0,0001

Research methodology The microsections for structure examination were subjected to grinding with sandpaper of 250 to 1200 granulation. Next, they were mechanically polished with the use of diamond pastes and chemically, in a solution of 90 ml CH3OH + 10 ml HNO3. For etching, two reagents were used with the following chemical compositions: –10 ml fluoric acid (48%) + 90 ml water; –2 ml fluoric acid + 2 ml nitric acid + 96 ml water. The observation was performed with: • REICHERT light microscope of MeF2 type; the results of the observation were recorded with a NIKON digital apparatus, COOLPIX 990 model; • HITACHI S-4200 scanning microscope with a cold cathode, equipped with an Xradiation detector EDS-VOYAGER of NORAN INSTRUMENTS. Research results As a result of the microscopic examination performed, it was found out that the casting AZ91D alloy is characterised by a solid solution structure α with α+β eutectic and β phase (Mg17Al12) at grain boundaries (Figs 4÷6). The eutectic obtained assumes the form of a so-called abnormal eutectic, where the initiating particles dissolved in the solid solution precipitate during the casting process at the boundaries of the initiating solid solution. Furthermore, the occurrence of Laves’ phase in the form of Mg2Si was proved as well as in the form of precipitations of probably a MnAl4 phase (Fig. 6). The Mg2Si phase has a light blue colour and is characterised by a regular, multilateral shape with smooth edges. The precipitations containing manganese and aluminium are characterised by an irregular shape with a rough surface; they very often take the form of spines. The closer the precipitation surface, the higher the ratio of Al to Mn content. The microanalysis results of the chemical composition of individual phases are shown in Table 2 (Figs 5, 6). a)

b) α+β

Mg2Si

Mg17Al12

α

Fig. 4. Microstructure of the AZ91 alloy cast: solid solution α + eutectic (α+β) + intermetallic phase Mg17Al12 and Mg2Si precipitations a) LM 150×; b) LM 400×;

194

a)

b)

c)

Fig. 5. Microstructure of AZ91 alloy. α + (α + β) + β a) SEM magn. 1500×; b) spectrum of α-solution chemical composition c) spectrum of β-phase chemical composition - Mg17Al12 Table 2. Chemical composition of identified AZ91 alloy phases

Phase

Chemical element [at.- %]

α phase β phase Mg2Si phase MnAl4 phase

Mg 89,41 63,39 62,39 1,37

Al 9,75 35,65 78,17

Zn 0,84 0,69 -

Si 37,61 -

Mn 20,47

Conclusions The research conducted was aimed at determining the microstructure of the casting AZ91D magnesium alloy which was cast into a sand mould. It has been shown that the analysed alloy has a solid solution structure α with an abnormal α + β eutectic formed as a result of precipitation, during casting, of β phase particles, with the β phase dissolved primarily in the solid solution; the alloy is also characterised by massive precipitations of β phase (Mg17Al12) at the solid solution grain boundaries. In addition, precipitations (light blue) of Mg2Si phase were observed, characterised by a regular, multilateral shape with smooth edges. Some acicular precipitations, of an uneven surface, of MnAl4 phase were found as well.

195

a)

b)

c)

d)

Fig. 6. Precipitations of Mg2Si and AlMn4 phases in the AZ91 alloy matrix a) SEM 500×, b) SEM 1000× c) spectrum of Mg2S phase chemical composition d) spectrum of AlMn4.phase chemical composition

References [1] L. Cizek et al.: ”Structure and properties of the selected magnesium alloys.”, 10th International Scientific Conference “Achievements in Mechanical & Materials Engineering”, Zakopane 2001. [2] A. Luo: ”Magnesium: current and potential automotive applications.”, Journal of Materials, February 2002. [3] Avedesian M., Baker H.: „Magnesium and Magnesium Alloys.”, ASM Speciality Handbook, 1999 [4] Dahle K., A., Lee Y. C., Nave M., D., Schaffer P., L., St Johnson D.H.: „Development of the as-cast microstructure in magnesium-aluminum alloys.”, Journal of Light Metals 1, 2001, 61-73

Microstructural Changes of AZ91 Magnesium Alloy After Heat Treatment Andrzej Kiełbus*, Lubomir Ciżek**, Libor Pawlica** * Silesian University of Technology, Katowice, Poland **VSB-TU Ostrava, Czech Republic

Abstract Mg-Al-Zn based alloys are the most commonly used magnesium alloys for structural components. AZ91 alloy is a widely used casting magnesium alloy. The paper presents microstructural characterisation of AZ91 (Mg-9Al.-0,6Zn) alloy after casting and heat treatment. The casting temperature was 740°C and heat treatment was performed at 415 °C/1÷48 h with aging at 165 °C for 8h. The microstructure of the casting alloy consists of α-Mg phase matrix with a primary brittle β phase (Mg17Al12) at grain boundaries. We also observed precipitations of Mg2Si and MnAl4 phases. The microstructural changes of this alloy have been investigated using light microscopy.

Introduction History and properties of magnesium alloys [1÷3] Magnesium alloys belong to the lightest structural alloys. They are characterised by low density: ~1,8 g/cm3, tensile strength Rm = 300÷350 MPa, elongation A5 = 20% and hardness ~100HB. Mainly for these reasons, magnesium alloys have a widespread application in the motor vehicle and aircraft industries as well as in household appliances and office machines. In the past, magnesium alloys played a significant role. During the First and Second World Wars they were used in the nuclear, metal and military aircraft industries. After the Second World War, the interest in magnesium alloys fell considerably. Only in the 90 s, about 17 kg components made of magnesium alloys were used for the production of VW Beetle. Requirements connected with the fumes emission limitation necessitated renewed growth of interest in those alloys. In 1994, magnesium consumption reached 228 thousand tons and in 1998, already 360 thousand tons, at 3.6 $ for one kilogram. It is forecast that within the nearest 10 years the consumption of magnesium alloys will grow at a 7% rate a year. As all materials, magnesium alloys have a number of advantages and disadvantages. Their advantages include first of all: low density, relatively high specific strength, good castability, very good for pressure die casting, good weldability in a controlled atmosphere,good corrosion resistance. The most important disadvantages are: low Young’s modulus, limited possibilities of cold workability, relatively low hardness, limited strength and creep resistance at elevated temperatures, great shrinkage during solidification, high chemical reactivity,in some cases, limited corrosion resistance. Improvement of the properties of magnesium alloys by conventional techniques (e.g. heat treatment) is not possible. The only solution is to apply reinforcing fibres, i.e. to produce a composite.

197 Heat treatment of magnesium alloys[1, 4] Magnesium alloys are subjected to heat treatment mostly for the purpose of improvement of their mechanical properties or as an intermediary operation, to prepare the alloy to other specific treatment processes. The type of heat treatment depends on the chemical composition of the alloy, its form (casting or after plastic working) and on the anticipated service parameters. Solution heat treatment of magnesium alloys enhances their strength, with maximum ductility and resistance to dynamic loads. Aging of solution heat treated alloys allows obtaining maximum hardness and yield point, with a decrease of ductility, whereas aging treatment without prior solution heat treatment, as well as annealing, result in a decrease of casting stresses and partial increase of mechanical properties when under tension. Annealing significantly decreases the mechanical properties and causes improvement of plastic properties, thus facilitating further treatment. A change of the heat treatment basic parameters has an influence on a change of the properties. For instance, extension of the annealing time considerably raises the yield point (with a decrease of ductility). The basic magnesium alloy heat treatment processes and their designations are presented in Table 1. Table 1. Designation of basic processes of magnesium alloys’ heat treatment

Designation O T2 T3 T4 T5 T6 T7 T8 T9 T10

Process Recrystallizing – only alloys for plastic working Annealing – only casting alloys Solution treatment and cold plastic working Solution treatment Aging treatment Solution treatment and aging Solution treatment and stress relief annealing Solution treatment, cold plastic working, aging Solution treatment, aging, cold plastic working Aging treatment, cold plastic working

The most frequently applied heat treatment processes for magnesium alloys are solution treatment and aging treatment. During solution treatment of Mg-Al-Zn alloys, the components subjected to treatment should be inserted in a furnace at ca. 260 °C and then, slowly heated with furnace to the solution treatment temperature to avoid melting the eutectic components and thus, to minimize voids formation. The heating rate from 260 °C to the solution treatment temperature depends on the size and chemical composition of the component subject to treatment as well as on its weight and differences in thickness of individual fragments of the cast. Most often, however the heating time is 2 hours. Some components, made of other magnesium alloys, can be inserted in the furnace at the solution treatment temperature. In the cases when the microstructure is incorrect or when a cast has been annealed too much during slow cooling after solution treatment, reheat treating is required. Majority of magnesium alloys subjected to reheat treating do not show an excessive grain growth. In the case of Mg-Al-Zn alloys, repeated solution treatment should not last longer than 30 minutes. After solution treatment, magnesium alloys are usually cooled in the open air. In some cases, however, other cooling agents are used, such as: water, glycol or oil. The two latter are mostly used to reduce casting stresses. The main problems which occur during and after heat treatment include: oxidation, formation of fusion voids, warping, grain growth

198 and incorrect properties. Oxidation occurs in the case of heat treatment without the application of neutral or protective atmosphere. It may lead to local weakening of the component treated, and even to burning of metal in the furnace. Voids are formed when inappropriate rate of heating from 260 to 370 °C is applied for Mg-Al-Zn alloys, as well as in consequence of exceeding the solution treatment temperature required for these alloys and for alloys with Zn and rare-earth elements additions. The phases situated at grain boundaries relocate along those boundaries and create long, narrow void regions. This process is usually connected with the growth of grain size. Warping occurs when the component subjected to treatment has no support or as a result of its non-uniform heating during heat treatment. The grain size growth usually occurs in AM100A, AZ81A, AZ91C and AZ92A alloys towards the end of solution treatment. The disadvantageous properties of a component subject to treatment occur as a result of too low or too high process temperature, bad circulation, too slow cooling after solution treatment or inappropriate time of holding thick components.

Experimental procedure Material for the research The object of the research was an AZ91D alloy designed for car wheel rims, cast into a sand mould. The chemical composition of the analysed alloy is shown in Table 2. Table 2. Chemical composition of the AZ91D alloy in wt-%.

Mg Rest

Al 9,15

Zn 0,6

Mn 0,24

Si 0,03

Cu 0,01

Fe 0,01

Be 0,0001

The casting technology included melting in a 250 kg induction furnace with the application of modification with SPEFINAL T 200 as well as refining with Emgesal Flux 200. The casting temperature was 740 °C [5]. After casting, heat treatment (solution and aging treatment) was performed. In order to protect the specimens against oxidation, the process was conducted in the atmosphere. The solution and aging treatment parameters are presented in Table 3. Table 3. Parameters of the heat treatment applied.

Designation O 1 2 3 4 5

Solution treatment Temperature Time As-cast 415 18 415 18 415 18 415 18 415 18

Cooling

Aging treatment Temperature Time

Cooling

air water furnace air water

168 168

air air

8 8

Research results As-cast alloy The casting AZ91D alloy is characterized by a solid solution structure α with eutectic α + β and β (Mg17Al12) phase at grain boundaries. Moreover, the occurrence of Laves’ phase in the form of Mg2Si and precipitations, probably of MnAl4 phase, has been proved (Fig. 1). The eutectic regions constitute ca. 24% and hardness of this material is 53HV.

199

a)

b)

Fig. 1. Microstructure of the as-cast AZ91 alloy – solid solution α + (α + β) eutectic + Mg17Al12 intermetallic phase and Mg2Si precipitations. LM 150×; b) LM 400×

Solution treatment 415 °C/18 h/air After solution treatment followed by air cooling, numerous precipitations were found in the alloy structure, of β (Mg17Al12) phase and MnAl4 phase in the α solid solution matrix. No occurrence of eutectic regions (Fig. 2) was detected. The alloy hardness (54 HV) did not change compared to the as-cast condition, and the mean area of the solid solution 2 grain was 1505 μm .

a)

b)

Fig. 2. Microstructure of the AZ91 alloy after solution treatment 415 °C/18h/air – a) LM 150×; b) LM 400×;

Solution treatment 415 °C/18 h/water After solution treatment followed by cooling in the water, similarly to air cooling, precipitations were found in the alloy structure of β (Mg17Al12) phase and MnAl4 phase in the α solid solution matrix (Fig. 3). No occurrence of eutectic regions was detected. The alloy hardness was the same as in the as-cast condition, i.e. 54HV. Solution treatment 415 °C/18 h/furnace After solution treatment followed by cooling in the furnace, the alloy is characterized by α solid solution structure with α + β eutectic regions; few precipitations of β (Mg17Al12) phase were observed at grain boundaries. Some precipitations of Mg2Si (Fig.4) phase were observed, as well. The structure of this material resembles that of the as-cast alloy, however, hardness after cooling with the furnace increases up to 65 HV; also, the number of eutectic regions increases (to ~36%) comparing to the as-cast condition.

200

a)

b)

Fig. 3. Microstructure of the AZ91 alloy after solution treatment 415 °C/18 h/water a) LM 150×; b) LM 400×;

a)

b)

Fig. 4. Microstructure of the AZ91 alloy after holding at 415 °C/18 h/furnace a) LM 150×; b) LM 400×;

Solution treatment 415 °C/18 h/air + aging 168 °C/18 h/air The aging treatment applied after solution treatment with air cooling caused β phase precipitation in the form of pseudoeutectic regions. Furthermore, the β phase (Mg17Al12) precipitated at solid solution grain boundaries (Fig. 5). Hardness of this material increased to 60 HV compared to that after solution treatment, whereas the solid solution area was simi2 lar and amounted 1532 μm .

a)

b)

Fig. 5. Microstructure of the AZ91 alloy after solution treatment 415 °C/18 h/air + aging 168 °C/18 h/air a)LM 150×; b) LM 400×;

201 Solution treatment 415 °C/18 h/water + aging 168 °C/18 h/air Aging treatment of an alloy solution treated and cooled in water caused precipitation of a larger number of β phase in the form of pseudoeutectic regions compared to the condition after air cooling (Fig. 6). Hardness of the AZ91 alloy after water solution treatment and aging is higher that that of the alloy after air solution treatment and aging, and amounts 65 HV, with the mean area reduced to 1065 μm2.

a)

b)

Fig. 6. Microstructure of the AZ91 alloy after solution treatment 415 °C/18 h/water+aging 168 °C/18h/air a)LM 150×; b) LM 400×;

Conclusions The subject of the research conducted was an evaluation of the influence of the heat treatment process on the microstructure and hardness of the AZ91 alloy. It was found that immediately after casting, an AZ91D alloy has an α solid solution structure with α+β eutectic and β (Mg17Al12) phase at grain boundaries. After solution treatment (both in the water and in the air), a reduction of pseudoetectic regions was observed as well as of the number of β phase precipitations. Hardness of the investigated as-cast and solution treated alloy remained unchanged, amounting 54 HV. Application of aging treatment caused precipitation of β phase along the solid solution grain boundaries and in the form of pseudoeutectic regions, which enhanced the increase of the alloy hardness.

References [1] Avedesian M., Baker H.: „Magnesium and Magnesium Alloys.”, ASM Speciality Handbook, 1999 [2] L. Cizek et al.: “Structure and properties of the selected magnesium alloys.”, 10th International Scientific Conference “Achievements in Mechanical & Materials Engineering”, Zakopane 2001. [3] A. Luo: “Magnesium: current and potential automotive applications.”, Journal of Materials, February 2002. [4] Stratton P.F., Chamg E.K.: “Protective atmospheres for the heat treatment of magnesium alloys. [5] Pawlica L., Čížek L., Greger M., Pavelcová J., Kiełbus A.: “Study of the formability of magnesium alloy AZ91.”, XI Sem.naukowe pt.: Nowe technologie i materiały w metalurgii i inżynierii materiałowej, Katowice, 2003.

Microstructure of Ultra Fine Grained Mg and Mg-10 wt.% Gd Prepared by High-Pressure Torsion J. Cizek1), I. Prochazka1), I. Stulikova1), 2), B. Smola1), 2), R. Kuzel1), V. Cherkaska1), R.K. Islamgaliev3), O. Kulyasova3) 1)

Charles University, Faculty of Mathematics and Physics, Prague, Czech Republic Zentrum für Funktionswerkstoffe GmbH, Clausthal-Zellerfeld, Germany 3) Institute of Physics of Advanced Materials, Ufa State Aviation Technical University, Ufa, Russia 2)

1

Introduction

Applications of Mg-based alloys at elevated temperatures are limited by the low melting point of Mg. This difficulty can be at least partially overcome by an addition of rare earth elements. It has been shown that Mg-based alloys with rare earth elements can be designed for operating temperatures above 300 °C [1, 2]. Mg-Gd represents one of the promising systems for novel Mg-based hardenable alloys with high creep resistance at elevated temperatures [3]. The precipitation of a metastable c-based centered orthorhombic β’ phase in Mg-Gd alloys with Gd concentration above 10 wt.% results in considerable age hardening [3]. Material hardening can be achieved also by decrease of the mean grain size. The grain boundaries represent obstacles for dislocation motion. It makes plastic deformation more difficult and increases material hardness. The yield stress is inversely proportional to the mean grain size through the well-known Hall-Petch relation. Recently it has been demonstrated that ultra fine grained (UFG) metals with grain size around 100 nm and no residual porosity can be produced by severe plastic deformation (SPD) [4]. A number of preparation techniques based on SPD have been developed so far. Among them, high-pressure torsion (HPT) and equal channel angular pressing (ECAP) are used most frequently. The smallest grain sizes have been achieved by HPT. The UFG materials prepared by HPT exhibit only weak texture and homogeneous structure, so that they are convenient for experimental studies. The UFG samples prepared by ECAP exhibit usually larger grain sizes than HPT made specimens. On the other hand, ECAP allows for a production of massive UFG samples, which makes it suitable for possible future industrial applications. A number of UFG metals exhibit favorable mechanical properties which consist in combination of very high strength and significant ductility as was demonstrated e.g. on HPT deformed Ti [5]. For this reason, it is highly interesting to examine microstructure and physical properties of UFG Mg-based light alloys. Following this purpose, microstructure investigations of UFG Mg and UFG Mg-Gd alloy were performed in the present work. The extraordinary properties of UFG materials are closely related with defects (grain boundaries, dislocations) introduced by SPD. Positron lifetime (PL) spectroscopy is a well-established non-destructive technique with high sensitivity to open volume defects [6]. It enables identification of the defect types present in the material studied and determination of defect densities [6]. Thus, PL spectroscopy represents an ideal tool for defect studies of UFG materials. It has been successfully employed in microstructure investigations of UFG Cu [7] and UFG Ni [8]. In the present work PL spectroscopy was combined with X-ray diffraction (XRD), microhardness measurements, and direct observations of microstructure by TEM.

203

2

Experimental

The samples of as-received technical purity Mg (99.95%) were investigated. In addition, another set of Mg samples was annealed at 280 °C for 30 min and then slowly cooled down inside the furnace. It led to the complete recovery of defects. Therefore, the annealed Mg was used as a defect-free reference sample. The Mg-10 wt.% Gd alloy was prepared from the technical purity Mg by squeeze casting. The as-cast material was homogenized at 500 °C for 6 hours. The homogenization annealing was finished by rapid quenching into water of room temperature. The UFG samples were prepared from the as-received technical purity Mg and the homogenized Mg-10 wt.%vGd alloy. The UFG microstructure was achieved by HPT [5]. The initial coarse-grained samples were deformed at room temperature by torsion up to the true logarithmic strain ε = 7 [5]. At the same time the samples were compressed with high pressure of 6 GPa. The UFG samples were disk shaped with diameter around 11 mm and thickness of 0.5 mm. A fast-fast spectrometer similar to that described in [9] was employed in the present work. 22 The timing resolution of the spectrometer was 170 ps (FWHM Na) at the coincidence count -1 rate 130 s . The PL spectra were decomposed into exponential components using a maximum likelihood procedure [10]. The PL measurements were accompanied by theoretical calculations of positron lifetimes using the atomic superposition method (ATSUP) [11]. Observations of microstructure were performed on the JEOL 2000 FX electron microscope operating at 200 kV. Thin foils for TEM were electropolished in a twin-jet device TENUPOL-2. X-ray studies were carried out with the aid of XRD7 and HZG4(Seifert-FPM) powder diffractometers using Cu Kα radiation. The XRD profiles were fitted with the Pearson VII function by the program DIFPATAN [12]. The microhardness was measured by the Vickers method at load of 100 g applied for 10 s using the LECO M-400-A hardness tester.

3

Results and Discussion

3.1

Coarse-Grained Samples

The reference Mg sample annealed at 280 °C for 30 min exhibits only single-component PL spectrum with lifetime τ1 = 225.3 ± 0.5 ps, see Table 1. It means the all positrons annihilate from the free state in this sample. The measured lifetime τ1 exhibits reasonable agreement with the theoretically calculated Mg bulk positron lifetime τb = 233.2 ps, see Table 2. Table 1. Lifetimes and corresponding relative intensities of the exponential components resolved in PL spectra of the studied samples. The errors (one standard deviations) are given in parentheses.

Sample

τ1 (ps)

I1 (%)

τ2 (ps)

I2 (%)

Reference Mg As-cast Mg Mg-10 wt.% Gd HPT Mg HPT Mg-10 wt.% Gd

225.3(5) 204(4) 220(4) 188(5) 210(3)

100 63(1) 90.9(6) 32.4(5) 34(2)

275(1) 301(2) 277(2) 256(3)

37(1) 9.1(7) 67.6(5) 66(2)

204 a)

b)

Fig.1: Bright-field TEM images of Mg: a) as-received sample b) HPT deformed.

A bright-field TEM image of the initial coarse-grained Mg specimen is shown in Fig. 1a. The mean grain size was found to be about 10 μm. Randomly distributed dislocations were observed inside the grains. The dislocation density ρ ≈ 5×1012 m-2 was estimated by TEM. The PL spectrum of this specimen is well fitted by two exponential components. Lifetimes and relative intensities of these components are listed in Table 1. The lifetime τ1 lies below the bulk positron lifetime for Mg. It means that it comes from free positrons. On the other hand, the second component with the lifetime τ2 represents a contribution of positrons trapped at defects. In view of the TEM observations, we consider that positrons are trapped at dislocations. The as-received Mg exhibits microhardness HV = 440 ± 40 MPa. Table 2. Positron lifetimes calculated using ATSUP method.

τ (ps) Mg – bulk Mg – monovacancy Mg – divacancy Mg – monovacancy associated with Gd atom

233.2 296.6 315.7 295.3

A bright-field TEM image of the homogenized Mg-10 wt.% Gd alloy is shown in Fig. 2a. Large coarse grains were observed. A low dislocation density below ≈ 1012 m-2 was found by TEM. The PL spectrum of this specimen is well fitted by two exponential components given in Table 1. The first component with the lifetime τ1 < τb can be attributed to free positrons, while the second one with the lifetime τ2 comes from positrons trapped at defects. The low dislocation density in the specimen approaches the lower sensitivity limit of PL spectroscopy [6]. Therefore, positrons trapped at dislocations cannot represent a noticeable contribution to PL spectrum. The second component with the lifetime τ2 represents a contribution of positrons trapped in quenched-in excess vacancies, which were ”frozen” in the sample due to the rapid quenching. This interpretation is supported by the lifetime τ2 ≈ 300 ps, which is remarkably higher than the lifetime 275 ps of po-

205 sitrons trapped in Mg-dislocations, and agrees well with the calculated lifetime of positrons trapped in Mg monovacancy given in Table 2. Monovacancies in pure Mg become mobile below room temperature [13]. Thus, free Mg monovacancies are not stable at room temperature and are quickly annealed out. Contrary to it, we observed that positrons are trapped in the monovacancies in the homogenized Mg-10 wt.% Gd sample. The explanation is that the monovacancies are bound to Gd atoms. The vacancy-Gd pairs are stable at room temperature. One can see from Table 2 that the lifetime of positrons trapped in Mg-monovacancy practically does not differ from the lifetime of positrons trapped in a vacancy-Gd pair. Thus, a free monovacancy and a monovacancy-Gd pair cannot be distinguished by PL spectroscopy. Nevertheless, an enhanced amount of Gd in the vicinity of vacancies was proved by coincidence Doppler broadening measurements [14]. Microhardness HV = 680 ± 50 MPa was measured on the homogenized Mg-10 wt.% Gdsample. It is in reasonable agreement with the hardness measured on this alloy after homogenization in [39]. a)

b)

Fig.2: Bright-field TEM image of Mg-10wt.%Gd alloy: a) homogenized 500 oC/6h b) HPT deformed.

3.2

UFG Specimens

A bright-field TEM image and an electron diffraction pattern of the HPT deformed Mg specimen are shown in Fig. 1b. Two different kinds of regions were observed: i) “deformed regions” with UFG grains (100-300 nm) and high dislocation density, and ii) “recovered regions” with substantially larger grains (≈ 1 μm) and practically free of dislocations. The “recovered regions” indicate a dynamic recovery of the microstructure during the HPT processing. However, the recovery is incomplete. The presence of the “deformed regions” with UFG grains is demonstrated also by the X-ray back-reflection pattern, which consists of continuous diffraction rings. The PL spectrum of the HPT deformed Mg was decomposed into two exponential components, see Table 1. The first component with the lifetime τ1 comes from free positrons. The lifetime τ2 of the second component corresponds well to the lifetime of positrons trapped at dislocations. Hence, in connection with the TEM results, we can conclude that positrons are trapped at dislocations inside the “deformed regions”. It should be pointed out that dislocations in the “deformed regions” are distributed relatively homogeneously, i.e. they are present in the grain interiors as well as close to the grain boundaries. It is an important difference of the HPT deformed Mg

206 from the HPT deformed Cu and Ni, where dislocations are piled up along grain boundaries, while the grain interiors remain almost dislocation-free [7,8]. The determined microhardness 600 ± 40 MPa is by 40% higher than that of the as-received Mg. A bright-field TEM image and an electron diffraction pattern for the HPT processed Mg-10 wt.% Gd alloy are shown in Fig. 2b. Contrary to the HPT deformed Mg, the Mg10 wt.% Gd has uniform UFG microstructure with the mean grain size about 100 nm. It indicates no dynamical recovery during the HPT processing. The electron diffraction pattern shows high-angle missorientation of neighboring grains. High density of homogeneously distributed dislocations was observed. The PL spectrum of the HPT deformed Mg-10 wt.% Gd alloy consists of two exponential components, see Table 1. The first component with the lifetime τ1 comes from free positrons. The lifetime τ2 of the second component is by ≈ 20 ps shorter than the lifetime of positrons trapped at Mg-dislocations. It indicates a slightly different nature of defects in the HPT deformed Mg-10 wt.% Gd. It is likely due to a segregation of Gd atoms along dislocations [14]. The specimen broke at 250-300 MPa during tensile test at room temperature without noticeable ductility. Only homogenized Mg-10 wt.% Gd, on the other hand, had ductility of 5% and the strength of 158 MPa [15]. The microhardness increases with the radial distance r being 1700 ± 300 MPa at the center of the specimen and 2300 ± 300 MPa at the margin. Similar behavior of microhardness HV was observed also in HPT deformed Cu [16] and it seems to be typical for a number of UFG metals prepared by HPT. HV reflects most probably the increase of strain with r. The mean value of HV of the HPT deformed Mg-10 wt. %Gd alloy is by 230% higher than that for the homogenized Mg-10 wt. % Gd sample. No microvoids were found in both HPT deformed specimens. It is another important difference compared to HPT deformed Cu [7] and Ni [8]. X-ray diffraction patterns of HPT deformed Mg and the Mg-10 wt.% Gd showed texture of the (00l) type which could be well described by the empirical formula Ihkl ~ Rhkl exp (-D sin 2ψ), with the angle ψ between the texture axis and particular reflection plane (Ihkl is measured integrated intensity of the corresponding diffraction profile and Rhkl is its theoretical intensity for the random grain distribution). The values of D = 5 and D = 2 for Mg and Mg-10 wt.% Gd specimens, respectively were found. For Mg-10 wt.% Gd alloy the lower broadening of (00l) profiles with respect to other peaks indicate dominating presence of 〈a〉 dislocations with Burgers vector b = 1/3.a. [ 2 1 10 ]. No significant broadening for pure Mg sample was detected so that the dislocation density should be less than about 1 × 1013 m-2. However, major contribution to the diffraction peaks comes from the “recovered regions” (see above) so that the “deformed regions” cannot be well characterized by XRD.

4

Conclusions

Microstructure of HPT deformed Mg and Mg-10 wt.% Gd was characterized. Two kinds of regions were found in the HPT deformed Mg: “recovered regions” with coarse grains and low dislocation density and “deformed regions” with UFG grains and high dislocation density. The HPT made Mg-10 wt.% Gd exhibits homogeneous UFG microstructure with high density of uniformly distributed dislocations and grain size about 100 nm. No micro-

207 voids were found. Both HPT deformed specimens exhibit significantly increased microhardness.

5

Acknowledgement

The work was financially supported by The Czech Grant Agency (contract 106/01/D049), The Grant Agency of Charles University (contract 187/2001), and by the DFG, Germany.

6

References

[1] G. Sigaudo in Proceedings of 3rd International Magnesium Conference (Ed.: G.W. Lorimer), The Institute of Materials, Cambridge 1997, p. 137. [2] D. Eliezer, E. Aghion F.H. Froes, in Magnesium 97 (Eds.: E. Aghion, D. Eliezer),Magnesium Research Institute, Ltd., Beersheva 1998, p. 343. [3] P. Vostry, B. Smola, I. Stulikova, F. von Buch, B.L. Mordike, phys. stat. sol. (a) 1999, 175, 491-500. [4] R.Z. Valiev, R.K. Islamgaliev, I.V. Alexandrov, Prog. Mat. Sci. 2000, 45, 103. [5] A.A. Popov, I.Y. Pyshmintsev, S.L. Demakov, A.G. Illarionov, T.C. Lowe, A.V. Sergeeva, R.Z. Valiev, Scripta Mater 1997, 37, 1089. [6] P. Hautojärvi, C. Corbel, in Proceedings of the International School of Physics “Enrico Fermi”, Course CXXV, (Ed. A. Dupasquier, A.P. Mills), IOS Press, Varena 1995, p. 491. [7] J. Cizek, I. Prochazka, M. Cieslar, R. Kuzel, J. Kuriplach, F. Chmelik, I. Stulikova, F. Becvar, R.K. Islamgaliev, Phys. Rev. B 2002, 65, 094106. [8] J. Cizek, I. Prochazka, M. Cieslar, I. Stulikova, F. Chmelik, R.K. Islamgaliev, phys. stat. sol. (a) 2002, 191, 391-408. [9] F. Becvar, J. Cizek, L. Lestak, I. Novotny, I. Prochazka, F. Sebesta, Nucl. Instr. Meth. A 2000, 443, 557. [10] I. Prochazka, I. Novotny, F. Becvar, Mat. Sci. Forum 1997, 255-257, 772-774. [11] M.J. Puska, R.N. Nieminen, J. Phys. F: Met. Phys. 1983, 13, 333-346. [12] R. Kuzel, DIFPATAN - program for powder pattern analysis, 1995 (http://www.xray.cz/priv/kuzel/difpatan). [13] M. Fahnle, B. Meyer, J. Mayer, J.S. Oehrens, G. Bester, in Diffusion Mechanisms in Crystalline Materials (Ed.: Y. Mishin et al.) MRS Symposia Proceedings No. 527, p. 23. [14] J. Cizek, I. Prochazka, F. Becvar, I. Stulikova, B. Smola, R. Kuzel, V. Cherkaska, R.K. Islamgaliev, O. Kulyasova, to be submitted. [15] F. von Buch, PhD. thesis, TU Clausthal, Clausthal-Zellerfeld, 1999, p. 186. [16] J. Cizek, I. Prochazka, G. Brauer, W.Anwand, R. Kuzel, M. Cieslar, R.K. Islamgaliev, phys. stat. sol. (a) 2003, 195, 335-349.

Development of New Grain Structure and Tensile Properties Improving in a Hot Pressed and ECAP Processed ZK60 Magnesium Alloy A. Galiyev, R. Kaibyshev, M. Almakaev, Institute for Metals Superplasticity Problems, Ufa

1

Introduction

Recently, it has been reported that fine-grained microstructure can be attained in Mg alloys by equal-channel angular pressing (ECAP) [1-3]. The mechanical properties of these Mg products could be significantly improved at room and elevated temperatures. For example, ECAP processed Mg alloys showed enhanced room temperature ductility and superplasticity at a low temperature and high strain rate [4-6]. To date, however, very limited data are available for relationship between ECAP processing conditions and the formed microstructures of Mg alloys as well as mechanical properties of ECAP processed specimens. No attempt was made to evaluate dependence of tensile properties of Mg alloys upon procedure of ECAP although it is well known that the microstructure and properties of materials produced by ECAP are dependent upon the processing route [7-10]. The present investigation was initiated to examine the influence of the processing route on the microstructure and mechanical properties of ZK60 Mg alloy.

2

Experimental

A ZK60 (Mg-5.8 %Zn-0.65 %Zr) alloy used in the present study was received as a commercially extruded bar of 90 mm in diameter. Bars for ECAP were machined from the extruded bar with diameters of 20 mm and lengths 80 mm. ECAP was conducted at 523K using a die having a channel angle of 90°. Samples were pressed repeatedly to 1, 2, 4 and 8 passes, using processing routes A, BC or C, in which the samples was rotated by 0°, 90° and 180°, respectively, in the same direction after each pass. The discs of 5 mm in thickness were sliced from the samples subjected to ECAP. Microstructures of the ECAP processed specimens were observed on the X plane where this plane was defined as lying perpendicular to the direction of pressing [11]. Tensile specimens were machined from the as pressed samples with guage sections of 2 × 3 × 6 mm3 and with the guage lengths lying parallel to the longitudinal axes. Tensile tests were carried out in air at room temperature at an initial strain rate 2.8×10-3 s-1 and elevated temperature 523 K with initial strain rates from 2.8 × 10-4 to 1.4 × 10-3 s-1.

209

3

Results and Discussion

3.1

Microstructures after ECAP

Typical microstructures are shown in Fig. 1 for the Mg alloy in the as extruded state and after ECAP, using the route A. In the as extruded state the microstructure was very inhomogeneous with the volume fraction of recrystallized grains of ~40% and the sizes of recrystallized grains scattered in the range of 2-20 μm. ECAP resulted in grain refinement and increase of recrystallized volume (Fig. 2a). The recrystallized grain size reduced significantly to 1-4 μm after four passes. It can be seen in Fig. 2a that recrystallization occurred more rapidly for routes BC and C, where the volume fraction of recrystallized grains increased to almost 90% after two passes, while for route A the recrystallized volume

Figure 1. Microstructures of (a) as extruded sample and samples after ECAP using route A for: (b) one pass, (c) two passes, (d, e) four passes and (f) eight passes.

210 80

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90

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80

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100

40

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80 60 Fine-grained fraction Coarse-grained fraction Total volume

40

2 passes 4 passes

20 0

(d) 0

90

180

Specimen Rotation [°]

Figure 2. Changes in recrystallized volume with processing route and number of pressings: (a) total recrystallized volume, (b) fraction of coarse-grained recrystallized structure, (c) fraction of fine-grained recrystallized structure, (d) effect of the processing route for two and four passes.

90% after two passes, while for route A the recrystallized volume of ~85% was attained only after four passes. Inspection of Figs. 1c-f showed that regions of fine grains with size of ~1 μm appeared in the recrystallized microstructure after two passes for all processing routes, however the volume of these regions was strongly depended on route used and number of pressings (Fig. 2c, d). This was the reason of decreasing the average grain size after two and four passes for routes BC and C in comparison with route A (Fig. 3). The average grain size after four passes was ~3.1, 1.6 and 2.1 μm for routes A, BC and C, respectively. It should be noted that route BC was more effective in attaining finer grain size. 3.2

Tensile properties at room temperature

Stress-strain curves of the extruded sample, 1-pass ECAP sample and 4-pass ECAP samples after processing through routes A, BC and C are shown in Fig. 4. The extruded sample exhibited higher 0.2% proof stress and ultimate tensile strength compared to the ECAP specimens. However, ECAP specimens exhibited much larger elongation. This is because of the fact that the basal plane is inclined to the extrusion direction in the ECAP specimen and as result the ductility was improved [4].

211 10

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300

200 as extruded 1 pass 4 passes - route A 4 passes - route BC 4 passes - route C

100

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0

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4

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Number of Pressings

0 0.0

8

Figure 3. Variation of the grain size as a function of number of pressings for samples pressed using routes A, BC and C.

0.1

0.2

0.3

Strain Figure 4. Stress-strain curves of as extruded sample and samples after ECAP using routes A, BC and C.

The values of the 0.2% proof stress, ultimate tensile strength and elongation to failure are summarized in Fig. 5. It can be seen that all three samples exhibited reasonably similar values of the 0.2% proof stress and ultimate tensile strength, but elongation to failure was higher after processing through route BC. It is concluded that a microstructure formed through route BC is more effective in increasing the elongation to failure. 3.3

Tensile properties at elevated temperature

Figure 6 shows the typical stress-strain curves at 523 K and an initial strain rate of 5.6 × 10-4 s-1 of the extruded sample, 1-pass ECAP sample and 4-pass ECAP samples after processing through routes A, BC and C. It is apparent from Fig. 6 that the extruded sample and 1-pass ECAP sample with coarser-grained microstructures exhibited much higher (a)

35

Elongation to Failure [%]

Tensile Properties [MPa]

360 320

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Route A Route BC Route C

240 0.2% Proof Stress

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10 5

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Figure 5. Variation of (a) the 0.2% proof stress, ultimate tensile strength and (b) elongation to failure as a function of number of pressings for samples pressed using routes A, BC and C.

212 70

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Figure 6. Stress-strain curves at 523 K and an 5.6 × 10-4 s-1 of as extruded sample and samples after ECAP using routes A, BC and C.

Figure 7. Variation of the strain rate sensitivity as a function of strain rate at 523 K for as extruded sample and samples pressed using routes A, BC and C.

stresses than the 4-pass ECAP samples with the fine-grained microstructure. The elongation increased with increasing number of pressings and it was due to grain refinement attained through ECAP. This trend was confirmed in Fig. 7, where the values of strain rate sensitivity, m, increased essentially from ~0.23 for as extruded sample to ~0.37 for the 1-pass ECAP sample and to the range of 0.42-0.46 for the 4-pass ECAP samples that was a direct consequence of grain refinement introduced by ECAP. As a result significant elongations were achieved in the 4-pass ECAP samples. On the other hand, the differences in the superplastic characteristics were found in samples processed by routes A, BC and C (Fig. 8). Three samples showed m values of ~0.42, 0.43 and 0.46 and elongations of ~480, 570 and 680%, respectively. This investigation demonstrates that the superplastic ductility was a maximum when using route C with a rotation of the sample by 180° between each pass through the ECAP die. This result is not

Strain Rate Sensitivity

0.40 0.35 0.30

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700

0.45

600 500 400 300 Route A Route BC Route C

200 100 0

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Figure 8. Variation of (a) the strain rate sensitivity at 523 K and (b) elongation to failure at 523 K and an -4 -1 5.6 × 10 s as a function of number of pressings for samples pressed using routes A, BC and C.

213 consistent with the earlier investigations for aluminum and aluminum-based alloys [9, 10] showing route BC was the optimum processing route for achieving maximum ductility. Apparently, peculiarities of the crystallographic slip of magnesium alloys lead to attainment of optimum microstructure and superplastic ductility when using processing route C. Further investigations are required for evaluation of the route influence on texture and properties of ECAP processed magnesium-based alloys.

5

Conclusions

(1) The grain size of an Mg-5.8%Zn-0.65%Zr alloy was reduced to ~3.1, 1.6 and 2.1 μm using the process of ECAP by routes A, BC and C, respectively, for 4 passes at 523 K. (2) Temperature room ductility was a maximum when using processing route BC in which the sample was rotated by 90° in the same sense between each consecutive pressing through the ECAP die. (3) Optimal superplastic ductility was achieved after pressing through route C with a rotation of the sample by 180° between each pass through the ECAP die: for example, elongations of ~480, 570 and 680% were obtained at 523 K and an initial strain rate of 5.6 × 10-4 s-1 in the samples processed by routes A, BC and C, respectively.

6

References

[1] Mabuchi, M.; Iwasaki, H.; Yanase, K.; Higashi K., Scripta Mater. 1997, 36, 681-686. [2] Mabuchi, M.; Ameyama, K.; Iwasaki, H.; Higashi K., Acta Mater. 1999, 47, 20472057. [3] Yamashita, A.; Horita, Z.; Langdon, T.G. Mater. Sci. Eng. 2001, A300, 142-147. [4] Yoshida, Y.; Cisar, L.; Kamado, S.; Kojima, Y., Mater. Trans., 2003, 44, 468-475. [5] Watanabe, H.; Mukai, T.; Ishikawa, K.; Higashi, K., Mater. Sci. Forum 2003, 419422, 557-562. [6] Matsubara, K.; Miyahara, Y.; Makii, K.; Horita, Z.; Langdon, T. G., Mater. Sci. Forum 2003, 419-422, 497-502. [7] Iwahashi, Y.; Horita, Z.; Nemoto, M.; Langdon, T. G., Acta Mater. 1997, 45, 47334741. [8] Ferrasse, S.; Segal, V. M.; Hartwig, K. T.; Goforth R. E., Metall. Mater. Trans. 1997, 28A, 1047-1057. [9] Oh-ishi, K.; Horita, Z.; Furukawa, M.; Nemoto, M.; Langdon, T. G., Metall. Mater. Trans. 1998, 29A, 2011-2013. [10] Komura, S.; Furukawa M.; Horita, Z.; Nemoto, M.; Langdon, T. G., Mater. Sci. Eng. 2001, A297, 111-118. [11] Furukawa, M.; Iwahashi, Y.; Horita, Z.; Nemoto, M.; Langdon, T. G., Mater. Sci. Eng. 1998, A257, 328-332.

Texture Formation and Texture Modelling of AZ31 Magnesium Wrought Alloy A. Styczynski [1], Ch. Hartig [1], R. Bormann [1], F. Kaiser [2], J. Bohlen [2], D. Letzig [2] [1] Arbeitsbereich Werkstoffphysik und –Technologie, Technische Universität Hamburg-Harburg, 21071 Hamburg, Germany [2] GKSS-Forschungszentrum Geesthacht GmbH, 21502 Geesthacht, Germany

1

Abstract

The texture formation during cold rolling of AZ31 magnesium wrought alloy has been investigated. Two initial microstructures – 1. rolled and recrystallized, 2. cast – were considered. A typical tilting of maxima by ± 15° in the (0002) pole figures was found to be characteristic for the texture after cold rolling of both examined microstructures. The measured texture evolution and plastic anisotropy (yield stresses and r-values) can be explained by a modified Taylor model assuming activity of the basal, prismatic and pyramidal slip systems and of the {0 1 12} < 0 1 11 > twinning system.

2

Introduction

The insufficient room temperature formability of magnesium wrought alloys is commonly ascribed to their anisotropic plastic behaviour [1, 2] and makes processing steps like rolling, extrusion etc. difficult, thus limiting their application for structural parts. The slip activities and critical resolved shear stresses (CRSS) in hcp magnesium have been reported for different slip and twinning systems [3-5]. According to these experimental data, the CRSS of the basal slip mode {0001} < 01 12 > is much lower than those of the non basal slip and twinning modes. It can be shown that a fulfilment of the von Mises condition requires at least the activation of one hard slip mode with c-component burgers vectors (i.e. pyramidal slip or {0 1 12} < 0 1 11 > twinning) [2,6]. The texture evolution and the plastic anisotropy, as obtained by experiments and model calculations, can be considered as a tool for the investigation of slip mode activity and for the verification of the CRSS’s in an anisotropic polycrystal. The origin of the cold rolling texture of hcp metals has been widely considered in literature [6-8]. It was assumed that for pure hcp metals the texture evolution depends strongly on the c/a ratio. For magnesium with a c/a = 1.622, near to the ideal value of 1.633, a basal texture was commonly observed [8]. The basal, prismatic and pyramidal slip modes with Burgers vector as well as the pyramidal slip mode with Burgers vector and also the twinning system {0 1 12} < 0 1 11 > are recognised in literature to be active deformation modes in pure magnesium [2-8]. However, this crude approach is not applicable in magnesium alloys. Due to variations in chemical composition, the strengths of slip and twinning systems are altered and hence changes in texture evolution occur.

215

3

Experimental Procedures and Results

Three samples of the AZ31 magnesium alloy were investigated corresponding to different initial microstructures: a commercial recrystallized sheet (A), a squeeze cast bar (B) and a commercial rolled sheet (C). From materials A and B specimens were prepared for cold rolling experiments. The dimensions of each sample amounted to 30 mm width × 6 mm thick × 50 mm length. The initial microstructures of the commercial sheet and of the squeeze cast bar are shown in Fig. 1a and Fig. 1b. The grain size was about 60 μm for the commercial recrystallized material (A) and about 350 μm for the squeeze cast bar (B). The volume fraction of the soft secondary phase Al17Mg12 amounted to less than 5% for both materials.

Fig. 1: Initial microstructures of commercial recrystallized sheet A (a) and squeeze cast bar B (b); microstructures after rolling: recrystallized sheet (c) and squeeze cast bar (d).

The rolling was carried out at room temperature on a two-high mill with a roll diameter of 200 mm. The thickness reduction for each step was fixed to a true strain ϕ = 0.1 (strain rate 0.7 s-1 < ϕ& < 1.5 s-1) and rolling was performed up to a total true strain ϕ = 0.51 for the recrystallized sheet and ϕ = 0.50 for the squeeze cast material, respectively, before the specimen began to fracture. The microstructures (Fig. 1c,d) after rolling show a high density of deformation twins, a typical feature of heavily cold deformed AZ31 alloy [9]. The textures of the rolled sheets were measured on a D8 DISCOVER X-ray Diffractometer (BRUKER AXS Inc.) using reflection geometry and Cu-Kα radiation. For each sample the {10 10}, {0002}, {10 11},{10 12},{1120},{10 13} and {1122} pole figures were recorded. The orientation distribution function (ODF) and complete pole figures were calculated using the harmonic method of Bunge [10]. The results of texture measurements are shown in Fig. 2.

216

Fig. 2: Initial textures (a: sheet A, b: cast material B) and textures after rolling (c: material A, d: material B)

The initial textures can be characterized as a strong basal texture for the recrystallized sheet and as a random texture with scattered, strong reflections from large grains for the squeeze cast bar (Fig. 2a,b). The texture after cold rolling can be completely characterized by a symmetrical tilting of the {0002} basal poles by ±15° towards the rolling direction (RD) (Fig. 2c,d). The tilting is more clearly visible for the cold rolled recrystallized sheet material than for the squeeze cast bar with initially nearly random texture. The {1120} pole figures give evidence for an almost perfect fibre texture: There are no distinct maxima visible at a polar angle α = 90° (Fig. 2cd). In terms of an ODF this texture can be described as two fibre components along the ϕ2 Euler angle (I.: ϕ1Φϕ2 = +90°,15°, ϕ2 ; II.: ϕ1Φϕ2 = −90°, 15°, ϕ2) being equivalent because of the orthorhombic rolling specimen symmetry. This fibre texture is clearly visible also in the squeeze cast material with a random initial texture. Thus it can be concluded that the basal pole tilting is merely caused by deformation mechanisms, not by the initial texture. For the measurement of the plastic anisotropy another typical commercial rolled sheet of AZ31 alloy (specimen C), 1.6 mm thick, and finally tempered under the H24 condition has been chosen for investigations. The texture of this rolled sheet was found to be almost equal with the texture of specimen A after cold rolling (c.f. Fig. 2c). The anisotropy parameter r was measured at room temperature by tensile tests of samples cut from the sheet at different angles β in the range between 0° and 90° to the rolling direction. Thereby the parameter r, defined as the ratio r = εW/εN, is obtained as a function of β, where εW is the plastic strain into the transverse direction of the tensile sample and εN the strain into the normal direction of the sheet. Results of this measurement, being presented in detail in [11], are compared in chapter 4 with theoretical calculations.

217

4

Texture Modelling and Plastic Anisotropy

Theoretical textures were calculated on the basis of a viscoplastic Taylor model [12] by usage of the LApp [13] software package. Two different initial textures: A random distribution, generated by a random generator, for the squeeze cast material and an orientation distribution, calculated from the experimental ODF before rolling, for the experimental recrystallized sheet were taken as input data. The strain rate sensitivity exponent m of the critical resolved shear stress (CRSS) was set to m = 19 for every slip system. The simulations were performed up to a total deformation of ϕ = 0.5 without strain hardening and with full Taylor constraints. A systematic analysis of possible texture types was performed taking into account all triplet

Fig. 3: Results of texture simulations for different combinations of slip and twinning systems (c.f. Table 1). A random initial texture was supposed and the total deformation grade was set to ϕ =0.5.

combinations of slip and twinning modes (c.f. Table 1) and a meaningful range of CRSS’s. Details of this analysis are described in [2] and representative results of the texture modelling are presented in Fig. 3 in form of {0002} pole figures. A typical maximum intensity splitting in the {0002} pole figures as in the experimental texture (Fig. 2) appears exclusively for the triplet combinations A, C and E (Fig. 3), allowing the slip mode, whereas it does not appear for the triplet combinations B and D, forbidding slip. This gives clear evidence that the activity of the pyramidal slip system is responsible for the basal splitting. The C combination (with the CRSS’s: τc( basal) = 1.0, τc( pyramidal2) = 3.0 and τc(twinning) = 2.0) is the only triplet combination, which does not exhibit any {0002} pole intensities along the rolling direction, as in the experimental textures. Thus, at a first glance this combination may be considered being active during cold rolling of AZ31. For a further verification of the mode combination C the theoretical yield locus and the rvalues of the tensile tested commercial sheet (c.f. chapter 3) were determined, taking into account the measured texture of this material. The same viscoplastic Taylor model as for the

218 Table 1: Combinations of slip and twinning systems used for theoretical texture calculations

designation A B C D E CE

slip and twinning combination b {0001} < 2 1 10 > + py1 {01 11} < 2 1 10 > + py2 {12 12} < 12 13 > b {0001} < 2 1 10 > +py1 {01 11} < 2 1 10 > + tw {01 12} < 01 1 1 > b {0001} < 2 1 10 > + py2 {12 12} < 12 13 > + tw {01 12} < 01 1 1 > b {0001} < 2 1 10 > + pr {01 10} < 2 1 10 > + tw {01 12} < 01 1 1 > b {0001} < 2 1 10 > + pr {01 10} < 2 1 10 > + py2 {12 12} < 12 13 > b {0001} < 2 1 10 > + pr + {01 10} < 2 1 10 > py2 + {12 12} < 12 13 > + tw {01 12} < 01 1 1 >

calculations of the theoretical texture (s.a.) was used, with identical viscoplastic parameters and CRSS’s . In Fig. 4a the yield locus for this case is shown, assuming a plane stress state (σ33 = 0). From this curve it must be concluded that combination C cannot fulfil the necessary experimental conditions for positive values of slopes on the yield locus curve at σRD = 0 and σTD = 0, which are required for positive values of the anisotropy parameter r. Such a behaviour is also expected because of the strong basal texture and the lack of slip modes accommodating plastic strains in the transverse direction.

Fig. 4 (a) Yield loci for commercial AZ31 sheet for slip and twinning combinations C and CE. (b) Comparison of the theoretical plane anisotropy parameter r (calculated by LApp [12]) with Experiment (r-values taken for 1% strain) as function of the angle β (∠ RD, Tensile direction).

Therefore the mode combination C was modified by addition of the prism slip mode (with τc( prismatic) = 3.6) resulting in mode combination CE. The yield locus for this case, also shown in Fig. 4a, gives a good accordance with experimental results: The slopes at σRD = 0 and σTD = 0 as well as the yield stresses in rolling and transverse direction (c.f. Table 2) agree reasonably well with experimental results.. Also the theoretical rolling texture for the mode combination CE fits with the measured textures (Fig. 3). Furthermore, the theoretical r-values as function of the azimuthal angle β coincident quite well with experimental measured values for a strain of 1%, as can be seen from Fig. 4b. It has to be pointed out that experimental r-values of this material can become much greater (r = 1.2-2.7 ) for a strain of 10% [11]. A reasonable cause for this behaviour is given by the anisotropic strain hardening of the different slip modes and may be explained within the scope of advanced texture models. The good agreement of textures and values for the plastic anisotropy (yield stresses and rvalues) obtained from theory with experimental results gives a strong evidence for the activity of the four slip and twinning modes: basal, prismatic, pyramidal (c + a)

219 ( {12 12} < 12 13 > ) and twinning ( {01 12} < 01 1 1 > ). This result was obtained by a calculation based on the Taylor model, giving some confidence to a “compatibility first” model. Table 2: Theoretical and Measured yield stresses for rolled commercial sheet C

Angle with rolling direction (Degrees)

Measured Yield Stress (MPa)

Theoretical Yield Stress for Combination CE

0 45 90

219 240 263

6.6 6.7 6.9

5

Conclusions

Theoretical texture predictions from a Taylor model give a good qualitative agreement with experiment. The results are quite similar with results obtained by more advanced models [6]. The yield locus and r-values of a commercial AZ31 sheet, as calculated after a viscoplastic Taylor model agree quite well with experimental values for low plastic strains obtained from tensile tests.

References [1] Kaiser F., Letzig D., Bohlen J., Styczynski A., Hartig Ch., Kainer K.U. Materials Science Forum 2003, 419-422, p. 315 [2] Styczynski A., Hartig Ch., Bohlen J., Letzig D., Scripta Mat., submitted [3] Kelley E. W., Hosford W. F., Transactions of AIME 1968, 24, p. 654 [4] Stohr J. F., Poirier J. P., Philos. Mag. 1972; 25: p. 1313 [5] Ando S., Tonda H., Materials Science Forum 2000, 350-351, p. 43 [6] Agnew S. R., Yoo M. H., Tomé C. N. Acta Mater 2001, 49, p. 4277 [7] Agnew S. R., Proc. Magnesium Technology 2002, ed. by H.I. Kaplan, TMS, 2002; 169 [8] Rollet A. D., Wright S. I., Texture and Anisotropy (Ed.: U. F. Kocks, C. N. Tomé , H. R. Wenk), Cambridge University Press, 1998, Chapter 5 [9] Lach E., Kainer K. U., Bohmann A., Scharf M., Magnesium Alloys and their Applications (Ed.: K. U. Kainer), Wiley-VCH, 2000, 354 [10] Bunge H.-J., Texture Analysis in Material Science, Cuvillier Verlag Göttingen; 1993 [11] Kaiser F., Letzig D., Bohlen J., Styczynski A., Hartig Ch., Kainer K. U., Materials Science Forum, 2003, 419-422, p. 315 [12] Aernoudt E., van Houtte P., Leffers T., Plastic Deformation and Fracture of Materials (Ed.: H. Mughrabi), VCH-Verlag, Weinheim, 1992, Chapter 3 [13] Kocks U. F., Kallend J. S., Wenk H.-R., Rollet A. D, Wright S. I.: Preferred Orientation Package – Los Alamos, 1994

Texture Development of AM20 Tensile Samples Under Load S.B. Yi1,2, H.-G. Brokmeier1,2, K.U. Kainer2, T. Lippmann2 1 2

Institute of Materials Engineering, Technical University Clausthal, Clausthal-Zellerfeld, Germany Institute of Materials Engineering, GKSS-Research Center, Geesthacht, Germany

1

Introduction

The potential applications of Mg and its alloys are rapidly increased especially in automobile industries. However, one of the great problems for industrial applications is the low formability and the high directional anisotropy of these alloys, which are caused by their crystallographic hexagonal structure. For filling up the von Mises criterion of 5 independent slip systems the activation of only basal slip is not sufficient, so that the profound investigations on easier and/or on harder conditions for the activation of other deformation systems are necessary for getting better mechanical properties. Since the crystallographic texture relates directly with the mechanical anisotropy and the activity of different deformation systems, many researchers are making effort to manifest the relationship between the texture and the deformation behavior under various loading conditions. The frequent difficulty in the interpretation of the texture development in Mg-alloys was the ambiguity in the combination of the active deformation modes [1, 2]. The present study on the in-situ texture variations during tension was carried out to contribute to a basic understanding of the relationship between the crystallographic texture and the activation of various deformation systems. Because texture reflects the deformation history and deformation mechanisms, the in-situ texture measurements give us meaningful information to the influence of initial texture on the activation of deformation systems and on the mechanical anisotropy of deformed materials. The great potential of synchrotron radiation is based on the high photon flux and the excellent brilliance. In the case of texture measurements a set of complete pole figures can be obtained in relatively short counting time when using an area detector. Thus synchrotron radiation is ideal for in-situ experiments. Hard X-rays have a penetration depth of 2.3 cm for Mg in the case of 100 keV photon energy. For realizing different initial textures, tensile samples were cut in three different directions from an extruded bar. The results of this study can be also applied to determine the texture variation of different industrial products because each tensile sample has the characteristic texture components which are often found in industrial products.

2

Experiments

The initial material was a rectangular extruded AM20 bar with a dimension of 10 × 90 mm2 in cross-section. The microstructure of the bar consists of equi-axed grains of 4.3 μm mean size. Round type tensile samples (DIN 50125) were cut in 0°, 45° and 90° to the extrusion direction to take into account the different initial textures, with 6 mm in diameter and 24 mm in gauge length. Figure 1 shows the (0002) pole figure of the initial bar and of the cutted tensile samples. As showed in figure 1 each tensile sample stands for a

221 L

(a)

(b)

(c)

(d)

T

Figure 1. The (0 0 0 2) pole figures of the rectangular extruded bar (a). The pole figure coordinate system is changed by cutting tensile samples. (b) stands for 0°, (c) for 45° and (d) for 90° tensile sample. The capital letters ‘L’ and ‘T’ on figure (a) indicate the extrusion and transverse direction, respectively. (intensity level = 1,2,3,4, Pmax = 4.99)

characteristic initial texture. The main loading axis of the samples corresponds always to the center of the pole figure. A universal testing machine (UTM), by which a sample can be loaded up to 20kN, was installed at beamline BW5, Hasylab at DESY (Germany) [3]. Figure 2 presents a schematic view of the beam line with the UTM set-up. The tensile test was carried out at room temperature with an initial strain rate of 6 × 10-4 sec-1. For pole figure measurements during loading the sample is rotated with the UTM around the omega-axis. The omega-rotation with this set-up can reach 145° because the loading shafts obstruct the incident and dif2 fracted beams. The monochromatic incident beam with a size of 1×1 mm and a wavelength of 0.12 Å was irradiated to the sample and the diffracted beam, Debye-Scherrer cone, was registered with an image plate area detector type MAR345. The image plate was positioned perpendicular to the beam. In order to get complete pole figures a set of image plate exposures were taken at every 5° in omega. Data evaluation for complete pole figures were taken from corresponded Debye-Scherrer cones after background correction using integral intensities. The orientation distribution functions (ODFs) were calculated using the iterative series expansion method with positivity correction. The (1 0 -1 0), (0 0 0 2), (1 0 -1 1) and (1 1 -2 0) pole figures were used for calculating the ODF to a degree given by Lmax = 19 [4]. 6

1

3

7

4 5

2

Axis for omega-rotation

1. Monochromator 2. Hard X-rays (0.12Å) 3. Incident slit (I) 4. Counter diode 5. Incident slit (II) 6. UTM with sample 7. Area detector

Figure 2. Schematic view of the beam line set-up with UTM installed in Hasy-Lab.

3

Results and Discussion

The flow curves during uniaxial tension show certainly that each sample has different deformation behavior (Fig. 3). The stress fall-offs on the flow curves coincided with pole

222 2

1

3

4 4

3

0°

90° 45°

2 1 1

2

3

4

Figure 3. The stress-strain curves of each specimen during in-situ loading, the stress fall-offs on curves correspond to the pole figure measuring positions. The marked points with numbers are present in this study.

figure measurements which were caused by mechanical problems. The texture measuring points at different strains correspond to the position after the yield phenomena (point 1), during strain hardening (point 2) and quasi-steady state of plastic deformation (points 3 and 4). 90° and 45° samples show low yield strengths of about 160 MPa which testify the activation of tensile twins and {0 0 0 2}<1 1 -2 0> basal slip with lower CRSS. The 90° sample shows higher strain hardening rate in the plastic deformation range than other samples. The 45° sample has the largest fracture strain at about 25.1 %. On the other hand, the specimen cut at 0° has the highest yield strength of 244 MPa because the activation of basal slip is restricted for this crystallographic direction. These specific mechanical behaviors can be explained by the different activities of deformation modes which relate directly with the initial texture of the tensile samples. Most of the basal planes were oriented perpendicular to loading axis in 0° sample, where Ini. (Fmax =11.6 ) P. 1 (Fmax =15.7 ) P. 2 (Fmax =16.4 ) P. 3 (Fmax =16.6 ) P. 4 (Fmax =16.6 ) o

ϕ2 = 0

ϕ2 = 30o

Figure 4. ϕ2 = 0o, 30o ODF-sections of the 0° sample, initial state and in-situ measured. The measuring points correspond to the numbers on the Fig. 3. (intensity level = 1,3,5,7,9,11,13,15)

223 Ini. (Fmax =12.1 ) P. 1 (Fmax =12.8 ) P. 2 (Fmax =10.9 ) P. 3 (Fmax =10.3 ) P. 4 (Fmax =12.0 ) o

ϕ2 = 0

ϕ2 = 30o

Figure 5. ϕ2 = 0°, 30° ODF-sections of the 90° sample, initial state and in-situ measured. The measuring points correspond to the numbers on the Fig. 3. (intensity level = 1,3,5,7,9,11)

the 45° sample has no basal plane perpendicular to the tensile axis but at 45° inclined from the tensile axis. Considering slip on the basal plane the Schmid’s factor of 45° sample is higher than in the case of 0°, therefore, this sample has the lowest yield strength. However the low yield strength of the 90° sample can not be explained by Schmid’s law. This sample had the basal planes mostly parallel to the loading axis, and this arrangement of polycrystals is very favorable for the activation of {1 0 -1 2}<1 0 -1 1> tensile twinning. Because the twin boundaries act as barrier against the sliding of slips during further deformation, this sample shows the highest stain hardening behavior. Figure 4 presents ϕ2 = 0° and ϕ2 = 30° ODF-sections of the initial 0° sample and those of in-situ measurements during tensile test. The initial texture is composed of {-1 2 -1 0} <0 0 0 1> (Euler angle {ϕ1, Φ, ϕ2} = 90, 90, 0) and a strong texture-fibre between {0 1 -1 0} <2 -1 -1 0> (0, 90, 30) and {0 1 -1 0}<0 0 0 1> (90, 90, 30). It should be mentioned that this type of texture is common for round or bar extruded Mg. After the yield phenomena the {-1 2 -1 0}<0 0 0 1> component in the initial sample disappears, and a slight increment of {0 1 -1 0}<2 -1 -1 0> component appears (point 2 in Figure 4). During further deformation this sample shows only a slight change of texture intensity. It is well known that the texture component in which <1 0 -1 0> is parallel to the loading axis is stable during uniaxial deformation [5]. The strengthening of <1 0 -1 0> parallel to the loading axis generally relates with the activation of -slip systems. The ODFs of the 90° sample are shown in figure 5. The texture of the initial 90° sample is for the most part composed of c-axis parallel and perpendicular to tensile direction, which is mainly found on industrially extruded tubes and rectangular bars [6]. Severe changes of the texture are observed after a small amount of plastic deformation, the texture components with c-axis parallel to tensile direction become rapidly weak and {0 1 -1 0}<2 -1 -1 0> component is newly developed. The fiber-component between {-1 2 -1 0}<0 0 0 1> and {0 0 0 1}<0 1 -1 0> has a maximum at {-1 2 -1 3}<1 -2 1 2> (90,45,0). This rapid change of the texture occurs by the activation of twinning, which impose large rotations to the crystals. By increasing deformation, the {0 1 -1 0}<2 -1 -1 0> component becomes stronger and the {1 2 -1 0}<0 0 0 1> component develops with a consumption of the fiber-component. Another change is the development of the {0 1 -1 0}<2 -1 -1 3> (60,90,30) component.

224 Since the 45° sample has a monoclinic symmetry the ODF, this sample is represented in the Euler space of ϕ1 = 0° ~ 360°, Φ = 0° ~ 90° and ϕ2 = 0° ~ 60°. Figure 6 shows the ϕ2 = 0° and 30° ODF-sections during tension. The initial texture of this sample is composed of two weak Initial ( Fmax = 12.3 )

Position. 1 (Fmax = 14.5 )

Position 2 (Fmax = 14.0 )

Position 3 (Fmax = 13.8 )

ϕ2 = 0o

ϕ2 = 30o

o

ϕ2 = 0

ϕ2 = 30o

Position 4 (Fmax = 13.8 ) o

ϕ2 = 0

ϕ2 = 30o

Figure 6. ϕ2 = 0 , 30 ODF-sections of 45° sample, initial state and in-situ measured. The measuring points correspond to the numbers on the Fig. 3. (intensity level = 1,3,5,7,9,11,13) o

o

components, {-1 2 -1 0}<0 0 0 1>, {0 1 -1 0}<2 -1 -1 0>, and two strong components, {-1 2 -1 2}<1 -2 1 -2> (315, 55 , 0), {0 1 -1 2}<0 1 -1 -1> (270, 45 , 30). This type of texture can be characterized by the basal planes and pyramidal planes which are inclined 45° from the loading axis. As referred to the flow curve of the 45° sample, this crystal-arrangement is very favorable for the activation of pyramidal -slip as well as basal -slip. During plastic deformation the basal planes, which were initially 45° inclined, has an extra rotation of totally about 12° in the direction of the perpendicular direction to loading axis. This rotation can be found on the ODF as the rotation of {-1 2 -1 2}<1 -2 1 -2> component to {-2 4 -2 3}<2 1 -3 2> (325, 65, 0). The pyramidal plane is also rotated about 10° in direction of the loading axis, which is described on the ODF as the rotation of {0 1 -1

225 2}<0 1 -1 -1> to {0 3 -3 4}<0 2 -2 3>. The strengthening of {0 1 -1 0}<2 -1 -1 0> (0, 90, 30) occurred by the consumption of {0 1 -1 2}<0 1 -1 -1> component, i.e. the intensity of {0 1 -1 2}<0 1 -1 -1> is decreased from 11.2 at initial state to 5.7 at measuring point 4, while the intensity of {0 1 -1 0}<2 -1 -1 0> is increased from 2.3 to 5.8. In view of the Schmid’s law, these final texture is more difficult for further -slip than the initial state. However the basal planes laid at about 57° to the loading axis until fracture act as the positive factor for sliding of -dislocations. Because of this factor the 45° sample shows highest fracture strain and a relatively low value of strain hardening. These experimental results from in-situ measurements can be used for the texture simulation which is necessary for more detailed study on the quantitative determination of activities of different deformation modes inclusive -slip.

4

Summary

In-situ texture measurements using hard X-rays with 100keV photon energy were carried out to examine the influence of the initial texture on the mechanical behavior and on the further development of texture in AM20. Results are summarized as follows: 1. By using hard X-rays and a UTM set up complete pole figures during loading could be taken successfully. 2. The results from the in-situ measurements present a variation of the texture that supports and gives fundamental information for texture simulation. 3. The initial texture has a strong influence on the mechanical properties of AM20. The 0° sample which has unfavorable texture components for -slip shows the highest yield o strength. The texture components of the 45 sample is favorable for the activation of -slips so that this sample has the lowest yield strength and the highest fracture strain. In case of the 90° sample the arrangement of basal planes perpendicular to tensile direction causes an activation of tensile twinning at the beginning deformation stage, so that a higher strength hardening and a dramatic change of texture occurs.

Acknowledgments This work has been funded by the German Ministry of Education and Research (BMBF) under the contract numbers 03BRE8CL and 05KS1MCA/2.

References [1] [2] [3] [4] [5] [6]

H. Friedrich and S. Schmann, Mat.-wiss. u. Werkstofftech. 2001, 32, 6 S. Zaefferer, Mater. Sci.& Eng. A 2003, 344, 20 H.-G. Brokmeier, U. Zink, T. Reinert and W. Murach, J. Appl. Cryst. 1996, 29, 501 M. Dahms and H.J. Bunge, J. Appl. Cryst. 1989, 22, 439 I.L. Dillamore and W.T. Roberts, Metall. Reviews 1965, 10, 271 R.A. Lebensohn, M.I. Gonzalez, C.N. Tome and A.A. Pochettino, J. Nuclear Mater. 1996, 229, 57

226

Strengthening Phases in Extruded Quasicrystal Containing Mg-Zn-Y Alloys Alok Singh*, A.P. Tsai*, M. Nakamura*, M. Watanabe** and A. Kato** * National Institute for Materials Science, Tsukuba 305-0047, Japan ** Toyota Motor Corporation, Toyota 471-8572, Japan

1

Introduction

There is a need for magnesium alloys with good strength at 473 K. A strengthening phase is required which provides microstructural stability at higher temperatures. A special phase, called the icosahedral quasicrystalline phase, exists in the alloy system Mg-Zn-RE, where RE is Y or a rare earth element. An icosahedral phase is unusual because it possesses fivefold symmetry and a quasiperiodic structure. Due to its quasiperiodic order, dislocations cannot move easily through it, which makes it hard and brittle. Therefore it can form effective composites with other metallic phase. In the Mg-Zn-Y system it occurs in equilibrium with the magnesium matrix [1– 3], showing a definite orientation relationship and with a coherent interface [4]. In dilute Mg-Zn-Y alloys, the icosahedral phase forms in intergranular spaces. This phase is stable up to the eutectic temperature, which occurs above 723 K. Thus it gives stability to the microstructure and retards grain growth. If a fine grain microstructure is obtained, its strength can be retained until temperatures of 473 K and above. A fine distribution of the icosahedral phase in Mg-Zn-Y alloys has been obtained by thermomechanical processing. Bae et al. [5, 6] have employed hot rolling at 673 K for magnesium alloys containing 2– 4.3 at% Zn and 0.2 – 0.7 at% Y to achieve good tensile strengths at room temperature (RT) (yield stress (0.2 offset, YS) of 220 MPa and ultimate tensile strength (UTS) of 370 MPa with 17.2 at% elongation). However, this level of strength is not retained at higher temperatures, though a high ductility is attained. We have employed extrusion as the thermomechanical means to obtain a fine microstructure of the icosahedral phase [7]. Alloy Mg95Zn4.2Y0.8 was extruded with a ratio 1 0 :1 at two different temperatures of 673 K and 523 K. Detailed microstructural study was performed with transmission electron microscopy (TEM) and correlated with tensile strength to determine the role of various phases in strengthening. Here we examine the effect of extrusion temperature and heat treatment on tensile strength.

2

Experimental Procedure

The alloy corresponding to composition Mg95Zn4.2Y0.8 was prepared by melting high purity elements in an electric resistance furnace. The alloy was homogenized by annealing at 673 K for 10 h and then extruded with a reduction ratio of 10 :1. One batch was extruded at 673 K while another at 523 K. Heat treatments were carried out by sealing the specimens under argon in silica tubes. Samples for TEM were prepared by ion milling. A

227 JEOL 2000FX-II microscope operated at 200 kV and a JEOL 4000EX-II microscope operated at 400 kV were used. Tensile tests were performed at a strain rate of 10– 3 /s on round specimens with a diameter of 4 mm and gauge length of 22 mm at RT, 473 K and 573 K on a Shimatzu Autograph machine.

Figure 1: Optical micrographs showing the grain structure and distribution of the icosahedral phase (dark) along the extrusion direction in the alloy extruded at (a) 673 K and (b) 523 K.

3

Results

3.1

Microstructure

The as-cast alloy had grain size of approximately 30 μm with the icosahedral phase at the grain boundaries. Extrusion at two different temperatures produced different grain sizes. After extrusion at 673 K the grain size was about 2 μm, while it was typically 10 μm on extrusion at 523 K, as shown in optical micrographs of Figure 1. The main feature of the microstructure was the icosahedral phase. In as cast alloy it occurred on grain boundaries. In extruded alloy it occurred as rows of intergranular particles of size about 500 nm, an example in Figure 2. The quasiperiodic nature of the icosahedral phase is brought out from this diffraction pattern. Pentagons are drawn connecting intense spots to bring out its fivefold symmetry. The diffraction spots are arranged not periodically, but quasiperiodically. Fine icosahedral phase particles were also distributed in the matrix. In alloys extruded at 673 K these particles occurred as nano-particles, well facetted on the basal and prismatic planes of the matrix shown in Fig 3a. The icosahedral phase showed a well-defined orientation relationship with the matrix. One of the icosahedral twofold axes occurred along the matrix hexagonal axis [4]. At the interface icosahedral twofold planes match basal and prismatic planes of the matrix. In samples extruded at 523 K, however, the icosahedral phase particles that occurred in the matrix were over 100 mm in size and rough shaped Figure 3b. Occasionally cubic W phase (Mg2Zn3Y3) was found to exist in the alloys [8].

228

Figure 2: An intergranular icosahedral phase particle after 673 K extrusion sample and its fivefold diffraction pattern.

Figure 3: (a) Nano-icosahedral particles in a grain of 673 K extruded sample. (b) Fragmented icosahedral phase particles distributed in the matrix after extrusion at 523 K. Fine precipitation of τ1 phase is also observed.

Complex precipitation of Mg-Zn binary as well as ternary phases occurred in the matrix. In the samples extruded at 673 K, the equilibrium phase MgZn occurred, apparently by nucleation on dislocations. In the samples extruded at 523 K, fine precipitation occurred in the matrix. These precipitates were mainly two kinds – fine rod-like precipitates, apparently of the precursor phase β1’ [9, 10], and another fine precipitate of a ternary phase τ1. Solutionizing was carried out at 673 K (below the eutectic temperature) to achieve finer precipitation in the matrix. Solutionizing at this temperature did not dissolve the icosahedral phase, but the icosahedral phase became rounded upon this treatment. Solutionizing produced a fine distribution of two phase – τ1 and tiny rods of β1’ (Figure 4). The phase τ1, from electron diffraction patterns, appears to be similar to a ternary phase τ, reported in Zn-Mg-Dy [11], where it was shown to be related to quasicrystalline structures. τ1 appears to have defect structures, and is considered to be a phase intermediate between Zn7Mg4 and τ. The τ1 precipitates exhibited a rough interface with the matrix and a strain around them. Upon further annealing at 473 K (10 h) the β1’ rods grew a few hundred nanometers in length. After tensile tests dislocation pile-up on these rods were observed. Microstructures after solutionizing are shown in Figure 5.

229

Figure 4: Precipitates after solutionizing. (a) β1’ rods along with an icosahedral phase particle in alloy extruded at 673 K. (b) τ1 phase precipitates in alloy extruded at 523 K.

Figure 5: Microstructure after solutionizing at 673 K. (a) In alloy extruded at 673 K. Two icosahedral particles are pinning the grain boundary. (b) Alloy extruded at 523 K. Interaction of dislocations with icosahedral phase particles on mechanical testing at RT. Arrow shows growth of τ1 phase over icosahedral phase. Its diffraction in inset.

3.2

Tensile Strength

Figure 6 shows the tensile strength of the extruded alloys at RT, 473 K and 573 K. Asextruded alloys showed good strength at room temperature. Extrusion at 673 K gave better strength at RT (YS 262 Mpa, UTS 332, elongation ~18%). However, the YS fell to less than half at 473 K, exhibiting no strain hardening and resulting in poor UTS. Upon solutionizing, however, substantial strain hardening was introduced, which raised the UTS from 125 Mpa to 170 MPa at 473 K, Figure 7a. Subsequent annealing at 473 K raised the YS to 145 MPa. The stress-strain curves are compared with those from a commercial extruded AZ61 alloy. A better microstructural stability of the Mg-Zn-Y alloy is indicated. In case of extrusion at 523 K, the as extruded alloy retained substantial strength at 473 K (YS 133 MPa, UTS 163 MPa). In this case too the elongation at 473 K is about 40%. The strength dropped to very low levels at 573 K in both cases, accompanied by nearly 100% elongation. Upon solutionizing, a substantial strain hardening was observed at 473 K resulting in a significant rise in the ultimate tensile strength (UTS), as observed in Figure 7b.

230

Figure 6: Tensile stress-strain curves at RT, 473 K for alloys extruded at (a) 673 K and (b) 523 K.

Figure 7: Stress-strain curves at 473 K and 573 K of alloys in as-extruded (extr.), solutionizing (soln.) and solutionized and annealed (ann.) condition. A curve for an extruded AZ61 alloy is shown for comparison.

4

Discussion

The primary strengthening in these alloys appear to be from the icosahedral phase. It forms strong interfaces with the matrix. Its occurrence at the grain boundaries provides strengthening and retards grain growth. It interacts strongly with grain boundaries and dislocation. An important point is that heat treatments at a high temperature of 673 K is possible without substantial grain growth even though the grain size is very fine in the extruded condition. No significant loss in strength at room temperature occurred after the solutionizing treatments. The higher temperature strength became better. Further strengthening is provided by precipitation within the grains. Presence of the τ1 precipitates seem to be responsible for word hardening, while lengthening of the β1’’ rods directly correlates with increase in YS. The solubility of yttrium decreases with lowering temperature and is almost nil at RT. The heat treatment and cooling rates are critical for nucleation of competing phases. Greater control is necessary in extrusion at 673 K. The extrusion at 523 K has several advantages. The icosahedral phase is distributed mor effectively in the alloy and phase equilibrium is not altered during the process, avoiding further heat treatment.

231

5

Summary

The microstructure and strength of an icosahedral quasicrystal containing Mg-Zn-Y alloy has been examined after extrusion at two different temperatures of 673 K and 523 K. The RT strengths after extrusion at these different temperatures were nearly the same; however, extrusion at lower temperature gave better higher temperature strengths in asextruded condition.

6

Acknowledgements

This work is partly supported by Japan Science and Technology Corporation.

7

References

[1] [2] [3] [4]

Z. Luo. S. Zhang, Y Tang, D. Zhao, Scripta Metall Mater. 1993, 28, 1513. A. Niikura, A.P. Tsai, A Inoue, T. Masumoto, Phil. Mag. Lett. 1994, 69, 35. A.P. Tsai, Y. Murakami, A. Niikura, Phil. Mag. A 2000, 80, 1043–1054. A. Singh, A.P. Tsai, M. Nakamura, M. Watanabe, A. Kato, Phil. Mag. Lett. (in press). D.H. Bae, S.H. Kim, D.H. Kim, W.T. Kim, Acta Mater. 2002, 50, 2343 –2356. D.H. Bae, M.H. Lee, K.T. Kim, W.T. Kim, D.H. Kim, J. Alloy Compounds 2002, 342 –445. A. Singh, M. Nakamura, M. Watanabe, A. Kato, A.P. Tsai, Scripta Mater. 2003, 49, 417 –422. A. Singh, A.P. Tsai, Scripta Mater. 2003, 49, 143 –148. J.B. Clark, Acta Metall. 1965, 13, 1281– 1289. L.Y. Wei, G.L. Dunlop, H. Westengen, Metall. Mater. Trans. A 1995, 26A: 1705 – 1716. E. Abe, A.P. Tsai, Acta Cryst. B 2000, 56, 915 – 917.

[5] [6] [7] [8] [9] [10] [11]

Grain Structure Characterization of AM60 Die Castings by Electron Backscatter Diffraction (EBSD) Measurements in SEM Hans I. Laukli, Otto Lohne, Lars Arnberg Norwegian University of Science and Technology, Trondheim

Introduction The fine-scaled microstructures in high pressure die cast (HPDC) magnesium alloys impede quantitative measurements of the grain size. The microstructure typically consists of large pre-solidified crystals [1] and fine grains that are too small to be easily resolved in the optical microscope. The grain boundaries are generally difficult to separate due to the inter-granular dispersion of divorced eutectic that consists of Al-segregated Mg and the eutectic β-Mg17Al12 phase [2]. Electron backscatter diffraction (EBSD) measurements in SEM are shown to be a more suitable method for investigations of the grain structure.

Experimental Box-shaped AM60 castings, with a wall thickness of 2.5mm, were produced at Norsk Hydro’s Magnesium Competence Center, Porsgrunn, in a Bühler SC42D 420 ton cold chamber HPDC machine (details in [1]). A complete AM60 die casting is displayed in Figure 1.

Figure 1. Complete AM60 die casting with biscuit, runners and overflows.

For investigations in the optical microscope the samples were ground, mechanically polished and etched in an acetic acid glycol solution [3]. For EBSD measurements in SEM the samples were ion milled in an Ion-Tech FAB 306 atom mill. A Jeol 840 SEM equipped with an EBSD detector was used for the grain size investigations with operation parameters as listed in Table 1. For instantaneous data processing, HKL Channel software was used [4].

233 Table 1. Ion milling and SEM* operation parameters.

Acceleration voltage [kV] 5 20*

Current Gun [mA] angle 5 15o -6* ~10

Time [h] ~4

Specimen Working distance [mm] tilt

Resolution/ step size [μm]

70o*

1*

~23*

The EBSD process When scanning the electron beam on the tilted specimen surface in the SEM, Figure 2 a), inelastic scattering occurs. Bragg diffracted electrons for particular planes are channeled differently which results in a change in intensity. Two cones for each crystallographic plane are formed and appear as straight bands, Figure 2 b), termed electron backscatter patterns (EBSP) [4]. Patterns from each grain are detected on a phosphorous screen and by a CCD-camera, and with EBSD acquisition-software the crystallographic orientation of grid points in the microstructure is determined. Qualitative EBSD measurements can be obtained in band contrast images, Figure 2 c). Dark gray appearance corresponds to a lower quality. This can result from overlapping patterns which is typically associated with grain boundaries. Lighter gray indicates a higher quality, e.g. in the center of a grain. The pattern quality can also be influenced by the surface condition of the specimen [5]. A non-uniform ion milling of the sample can result in non-indexed points as observed in the white area, marked 1, in Figure 2 d). Occasionally, grid points are misindexed due to poor pattern quality, and appear as isolated and highly misoriented pixels on the orientation map [5], e.g. the pixels marked 2 in Figure 2 d). Different crystallographic orientations combined with grain boundaries can be processed to appear as grain maps, where a specific colour indicates a grain and grain boundaries are depicted as black lines. Misorientations greater than a specific value, e.g. 15°, are defined as grain boundaries.

Figure 2. a) Secondary electron image of the ion milled specimen surface, b) EBSP of a magnesium grain, c) band contrast image and d) orientation map (resolution 0.5 μm step size).

234

Results In Figure 3 a) an optical micrograph of the HPDC AM60 microstructure is displayed. A few coarse crystals is revealed by the etching, while the fine grains are difficult to resolve. An orientation map obtained with EBSD in SEM from an equivalent microstructure is observed in Figure 3 b). The average grain size is measured to be approximately 7 μm.

Figure 3. a) Optical micrograph of AM60 HPDC microstructure revealed by etching and b) orientation map of equivalent microstructure with an average grain size of 7μm.

Grain maps from adjacent areas can be stitched together to form quantitative orientation maps of the grain structure in large areas of the microstructure. In Figure 4 (bottom image) the grain structure is revealed by means of orientation maps approximately halfway through the thickness of the micrograph in Figure 4 (top image). Relatively coarse grains are present close to the surface and finer grains are observed at an intermediate position. In the center, coarse grains are present surrounded by fine grains.

Figure 4. Optical micrograph of etched cross section (top) and stitched orientation maps from cross section (bottom).

235

Acknowledgements The present work was funded by the project NorLight Shaped Castings and partners: the Norwegian Research Council, Alcoa Automotive Castings, Scandinavian Casting Center ANS, Elkem Aluminium ANS, Fundamus AS, Hydro Aluminium Metal Products, Norsk Hydro ASA, NIMR, NTNU, and SINTEF (project responsible). The authors thank the partners for the financial support.

References [1] H.I. Laukli, O. Lohne, S. Sannes, H. Gjestland, L. Arnberg, Int. J. C. Met. R., 2003. Submitted. [2] A.K. Dahle, Y.C. Lee, M.D. Nave, P.L. Schaffer, D.H. StJohn, J. Light Met., 2001, 1, 61-72. [3] Microstructure of magnesium based diecasting alloys in Hydro Diecaster Bulletin, 3, 1999. [4] HKL Technology Channel, 5, HKL Technology, 2001. [5] F.J. Humphreys, J. Mat. Sci., 2001, 36, 3833-3854.

Possibilities of ECAP of Magnesium Alloy Katarzyna N. Braszczyńska Institute of Materials Engineering, Technical University of Cz stochowa, stochowa, Al. Al. Armii Krajowej 19, 42-200 Cz st stochowa, Poland

Abstract The development of microstructure during equal-channel angular pressing (ECAP) of commercial AZ91 alloy was investigated. In order to obtain high plastic deformation structure warm die with the channel angle Φ and the corner angle Ψ equal 90° and 120°, respectively, was used. Samples were deformed isothermally at different temperatures and at the speed of 0.83 mm/s. The obtained results demonstrate the possibility of obtaining the homogeneous microstructure, characterized by considerable grain refinement in the investigated alloy.

Introduction Several plastic deformation (SPD) techniques are known as a tool to produce fine-grained materials, which can exhibit superplastic behavior in a certain range of thermo-mechanical parameters [1-7]. Although several different procedures are currently available for imposing the severe strains necessary to attain an ultrafine-grained structure, most attention has been devoted to the two techniques of high-pressure torsion (HTP) and equal channel angular pressing (ECAP). Processing through the procedure of ECAP is a metal working process allowing to obtain ultra-fine grain in bulk materials [6-16]. This procedure differs from more conventional processes, such as rolling and extrusion; it has the ability to introduce on exceptionally high plastic strain without any concomitant change in the cross-sectional dimensions of the sample. Equal-channel angular pressing (ECAP) in which a piece of material is pressed to pass two connected channels having equal cross sections and intersecting at an angle has emerged as a promising process to obtain new structure and properties of materials [7, 12, 17-23]. In practice, the strain imposed on the sample in ECAP depends upon the two angles defined in Fig. 1: there is an angle Φ between the two channels and an additional angle Ψ delineating the outer arc of curvature at the point of intersection of the channels. The microstructure and properties of materials pressed by ECAP are strongly dependent on the plastic deformation behavior during pressing which is governed mainly by die geometry, and process variables as temperature, pressing speed and rotation of the billet around its longitudinal axis between adjacent passes. Theoretically, the equivalent strain, ε, generated in the work piece after one pass of ECAP is given by the following relation [6]: ε=

1 ⎡ ⎛Ψ Φ⎞ ⎛ Ψ Φ ⎞⎤ ⎢ 2ctg ⎜ 2 + 2 ⎟ + Ψ cos ec ⎜ 2 + 2 ⎟ ⎥ 3⎣ ⎝ ⎠ ⎝ ⎠⎦

(1)

237

Figure 1. A schematic representation of ECAP die with the channel angle Φ and the corner angle Ψ.

Recently many workers are interested in applying ECAP process to magnesium alloys which are the lightest structural metals and hence, have many potential applications [1623]. However, magnesium and its alloys generally exhibit poor workability to improve because of HCP structure, which has limited slip systems resulting in the poor ductility. Grain refinement during the ECAP of metals with cubic structures is controlled by the formation of sub-boundaries with high dislocation densities. By contrast, the deformation of magnesium may be expected to be quite different because it involves not only the slip of dislocations but also deformation twinning. In the present study, equal channel angular pressing was undertaken in order to examine possibilities of grain refined in an AZ91 magnesium alloy.

Experimental material and procedure As-received AZ91 magnesium alloy had a dendritic microstructure, typical for the gravity-cast alloy, which was characterized by very heavy segregation of alloying elements. Mg-Al alloys are prone to segregation due to relatively wide temperature spans between the liquids and the solids curves. Non-equilibrium solidification conditions caused the formation of large crystals of the (Mg) phase (depleted in aluminum) and pushing the Al admixture away into inter-dendrical spaces (Fig. 2). The details of microstructure were described in work [24]. Before the ECAP processes, the samples were heat treated in order to obtain homogenizing microstructure, characterized by higher workability. The solution annealing was carried out at 693 K for 26 h in protective argon atmosphere. The microstructure obtained after this heat treatment is shown at Fig. 3. It should be noted that this homogenizing treatment is one of main processes for magnesium alloys, however, it causes obtaining great grains. The investigated AZ91 alloy was cut into rods with about 50 mm length and 11.8 mm diameter. The ECAP die used in this investigation was designed to obtain a shear strain of ~1.15 during each pass. It contained an inner contact angle Φ equal 90° and the cor-

238 ner angle Ψ of 1209. All billets were processed at a processing rate of 0.83 mm/s and by route BC. Molybdenum disulphide (MoS2) was used as lubricant.

Figure 2. Microstructure of AZ91 as-received

Figure 3. Microstructure of AZ91 alloy obtained after solution annealing

The processing temperatures were chosen for the best combination of ductility and efficient grain refinement. For each separate pressing, the sample was sprayed with lubricant and it was then placed in the channel after the die had reached the required temperature. Before the first press, sample was heated to a required temperature. Following ECAP, microstructure was observed using specimens cut either parallel or normal to the longitudinal axis.

Results Fig. 4. shows the macroscopic images of samples, which experienced a single ECAP pass under isothermal conditions at three different temperatures. At temperature below 553 K investigated alloy exhibited failure characterized by the formation of a series of segments along the length in one pass ECAP process (Fig. 4a, b). As temperature was raised, a transition in flow from segmented to uniform samples was observed (Fig. 4c). It should be noted that the billet temperature during an ECAP operation may be above the die temperature due to the heat generated by the mechanical work. As was analyzed [20], the maximum temperature rise at high strain rates was estimated as 26 K for AZ91 alloy.

239

a)

b)

c) Figure 4. Macrographs of AZ91 samples processed via ECAP at 473 K (a), 523K (b) and 553K (c); one pass. Direction of exit channel was marked with an arrow.

Representative microstructures in AZ91 ECAP samples pressed at different conditions are shown in Fig.5-7. Fig. 5a was taken from the transverse section of the as-pressed billet while Fig. 5b was taken from the longitudinal section sample passed at 553 K. In the transverse section, the small grains surrounding some larger grains were observed. In longitudinal section, the formation of a banded structure lay essentially parallel to the shearing plane at approximately 45° to the top and bottom surface of the sample. a)

b)

Figure 5. The microstructure of AZ91 after ECAP process at 553 K; two passes; the transverse section (a) and the longitudinal section (b)

Fig. 6 shows that same evidence for a retention of the banded structure after two passes through the die at 623K; the microstructure is homogeneous consisting equiaxed grains. Fig. 7 shows that after four passes at this temperature a reasonably homogeneous microstructure of small and equiaxed grains exists. The changes of shape and size of grains shown on these figures can appear as a consequence of ‘rotation dynamic recrystalization’ (RDX) during ECAP process. This dynamic recrystallization mechanism, described in works [25-26], is different from conventional

240 dynamic recrystallization (DRX) and was observed in magnesium alloys after hot compression or hot rolling processes. It is also apparent that ECAP is especially effective in reducing the grain size when the pressing is conducted at lower temperatures, where the occurrence of grain growth is limited. However, the grain size did not change during the ECAP processes carried out at temperatures below 673 K. This indicates, that the dynamic recrystallization plays main role in the grain refinement. a)

b)

Figure 6. The microstructure of AZ91 after ECAP process at 623 K; two passes; the transverse section (a) and the longitudinal section (b)

a)

b)

Figure 7. The microstructure of AZ91 after ECAP process at 623 K; four passes; the transverse section (a) and the longitudinal section (b)

Summary 1. The ECAP process was successfully carried out isothermally for AZ91 alloy in order to obtain grain refinement. 2. At temperatures below 553 K, the ECAP process carrying out in external shape of die was practically impossible, because of destruction of samples. 3. After pressing at temperature above 553 K the grain size of AZ91 alloy was drastically decreased and above 623 K homogeneous microstructures were obtained.

241

Acknowledgements This work was supported by State Committee for Scientific Research (KBN), Poland under grant 7T08B00222. The author is grateful for the fellowship from the Foundation for Polish Science.

References [1] M. Kumachi, M. Furukawa, Z. Horita, T.G. Longdon, Mater. Sc. and Eng., A347, 2003, 223-230 [2] C. Xu, T.G. Longdon, Scrip. Mater., 48, 2003, 1-4 [3] H.S. Kim, M.H. Seo, S.I. Hong, J. of Mater. Proc. and Tech., 130-131, 2002, 497503 [4] S.J. Oh, S.B. Kang, Mater. Sc. and Eng., A343, 2003, 107-115 [5] M. Furukawa, Z. Horita, T.G. Longdon, Mater. Sc. and Eng., A332, 2002, 97-109 [6] Y. Iwashashi, J. Wang, Z. Horita, M. Nemoto, T.G. Longdon, Scrip. Mater., 35, 1996, 143-146 [7] H.S. Kim, M.H. Seo, S.I. Hong J, J. of Mater. Proc. and Tech., 113, 2001, 622-626 [8] V. Stalyarov, Y.T. Zhu, I.V. Alexandrov, T.C. Lowe, R.Z. Valiev, Mater. Sc. and Eng., A299, 2001, 59-67 [9] I. Kim, J. kim, D.H. Shin, C.S. lee, S.K. Hwang, Mater. Sc. and Eng. A342, 2003, 302-310 [10] S.L. Semiatin, D.P. DeLo, Mater. and Des., 21, 2000, 311-322 [11] Z. Horita, T. Fujinami, T.G. Langdon, Mater. Sc. and Eng., A318, 2001, 34-41 [12] S. Komura, M. Furukawa, Z. Horita, M. Nemoto, T.G. Langdon, Mater. Sc. and Eng., A297, 2001, 111-118 [13] D.h. Shin, I. Kim, J. kim, Y.S. Kim, S.L Semiatin, Acta Mater., 51, 2003, 983-996 [14] Y. Nishida, H. Arima, J.C. Kim, T. Ando, Scrip. Mater., 45 2001, 261-266 [15] Y. Wu, I. Baker, Scrip. Mater., 37, 1997, 437-442 [16] T. Mukai, M. Yamanoi, H. Watanabe, K. Higashi, Scrip. Mater., 45, 2001, 89-94 [17] W.J. Kim, C.W. An, Y.S. Kim, S.I. Hong, Scrip. Mater., 47, 2002, 39-44 [18] M. Mabuchi, H. Iwasaki, K. Yanase, K. Higashi, Scrip. Mater., 36, 1997, 691-686 [19] A. Yamashita, Z. Horita, T.G. Langdon, Mater. Sc. and Eng., A300, 2001, 142-147 [20] Y. Nishida, T. Ando, M. Nagase, S. Lim, I. Shigematsu, A. Watazu, Scrip. Mater., 46, 2002, 211-216 [21] H. Watanabe, T. Mukai, K. Ishikawa, M. Mabuchi, K. Higashi, Mater. Sc. and Eng., A307, 2001, 119-128 [22] M. Mabuchi, K. Ameyama, H. Iwasaki, K. Higashi, Acta Mater., 47, 1999, 20472057 [23] W.J. Kim, C.W. An, Y.S. Kim, S.I. Hong, Scrip. Mater., 47, 2002, 39-44 [24] Braszczyńska K.N., Zeitschrift für Metallkunde, 93, 2002, 845-850 [25] S.E. Ion, F.J. Humphreys, S.H. White, Acta Metall., 30, 1982, 1909-1914 [26] J.A. Valle, M.T. Perez-Prado, O.A. Ruano, Mater. Sc. And Eng., A35, 2003, 68-78

Electron Probe Micro Analysis of Sedimented Zirconium Particles in Magnesium Bruce Davis* and Keyna O’Reilly** *Materials Science Department, Oxford University, now Magnesium Elektron, Manchester **Materials Science Department, Oxford UniversityJohn King, Magnesium Elektron, Manchester

Introduction The grain refinement of Mg and its alloys by Zr was first achieved in 1947 by Sauerwald [1]. Zr is a potent nucleant of Mg and can result in grain sizes of the order of 20μm in some alloy systems, a factor of ten smaller than can be achieved in the grain refinement of Al alloys and better than any other refinement system available for Mg alloys. The high potency of Zr particles as nucleants for solid Mg is due to: the small size of the particles (≈1μm); the close match of the lattice parameters and similarity to the crystal structure of Mg; and the process by which Zr particles are evenly dispersed throughout the melt. The equilibrium phase diagram for the Mg-Zr binary alloy system contains a peritectic reaction at 654 °C for Zr concentrations in excess of 0.58wt%. The process of grain refinement is thought to depend on particles of α-Zr that precipitate from solution as the melt cools below the liquidus temperature. This results in a cascade of very fine particles forming throughout the melt that act as nucleants of solid Mg below the peritectic temperature. It has been stated in the literature that grain refinement occurs in Mg even when the levels of Zr are below 0.58wt% [2]. This would suggest that undissolved Zr particles, introduced from the hardener play a role in grain refinement of Mg. However, it is generally considered that un-dissolved Zr does not act as a nucleant of Mg due to some difference of the particle surface, such as an oxide film, that interferes with the chemical wetting properties [3]. This study investigates the use of sedimentation as a method of separating Zr particles that formed in molten Mg for further analysis. Electron probe micro analysis was used to characterise the chemical distributions of particles within the samples. The results obtained are presented to help explain the phenomena of grain refinement in Mg containing low levels of Zr.

Experimental Methods and Materials The technique of sedimentation utilizes the difference between the density of intermetallic compounds present in the molten metal to separate them from the bulk of the liquid by allowing the particles to settle to the bottom of the sample. It is then possible to section the layer of sediment from the material after solidification for further analysis. Sedimentation has previously been used to study the effect of undercooling on the nucleation of TiAl3 in aluminium alloys [4, 5].

243 The sedimentation process was carried out in a three-zone vertical tube furnace for accurate temperature control and to minimise the temperature gradient along the length of the samples. This was measured to be less than 1°C. Accurate temperature control was required to prevent undershooting of the temperature when the samples were cooled. The furnace was mounted on a rubber base to minimise vibration, as this is a known cause of nucleation in solidifying metals. A dynamic argon atmosphere prevented oxidation of the liquid metal during experiments. All experiments were carried out in triplicate. The sedimentation studies were carried out on the grain refinement of commercially pure magnesium. Zirconium was added to the melt in the form of a proprietary hardener called Zirmax® supplied by Magnesium Elektron. Samples for use in the sedimentation experiments were grain refined using 5wt% Zirmax®, then cast for use in the sedimentation furnace in a mild steel die. The grain refinement and casting of the samples for sedimentation was carried out in accordance with Magnesium Elektron’s recommended practice. The melt was vigorously stirred following the addition of the Zirmax® and held for ten minutes prior to casting. The casting temperature was 780 °C.

Experimental Procedure for the Sedimentation Technique. Stage one of the sedimentation experiment involved cleaning the metal at the casting temperature, 780 °C. This removed any undissolved Zr particles that may have been present in the melt. When the crucible holder reached the required temperature it was raised to the top of the furnace. Three samples, contained in graphite crucibles, were placed in the holder and lowered into the hot region of the furnace. The samples were allowed to equilibrate for 45 minutes and then held for one hour to allow any particles present to settle to the bottom of the graphite crucibles. The crucible holder was then raised to the top of the furnace and the samples placed in a quench block. The quench block was a water-cooled brass crucible holder that chilled the base of the samples where the sedimented layer was present. When fully cooled the base of the samples were cut off, removing the sedimented layer. The remaining lengths of samples where then placed back in the crucibles for the next stage of the sedimentation experiment. Stage two of the sedimentation experiments involved reheating the samples to the original sedimentation temperature, 780 °C, to dissolve any Zr particles present in the bulk of the sample. The samples were then cooled by 20 °C and held for one hour. This allowed any Zr particles formed during cooling to settle to the base of the samples. The samples were then removed from the furnace and quenched as outlined above, and the sedimented layer removed. The process was repeated, cooling the samples by a further 20 °C each time, to produce material that had been held at 760, 740, 720, 700 °C. The grain size of these samples was measured using the linear line intercept method. A control experiment was performed to establish whether sedimentation of a liquid sample resulted in all solid particles present in the melt being removed. This was achieved by cleaning a sample through sedimentation at 780 °C as in stage one, removing all heavy inclusions, i.e. undissolved Zr particles, and quenching. The sedimented layer was then removed. The sample was reheated to 780 °C and held for one hour to allow any remaining or newly introduced particles to sediment to the base of the sample. The sample was then quenched and the base of the sample removed. Finally the sample was reheated to 780 °C and cooled to 720 °C, held of one hour, then quenched. The sections containing sedimented layers were inspected using optical microscopy.

244

Electron Probe Micro Analysis, EPMA EPMA of the samples was carried using a JEOL JXA8800 microprobe. The apparatus utilised wavelength dispersive x-ray analysis with four spectrometers, allowing four elements to be detected at any one time. The accelerating voltage used was 20 keV, with a dwell time of 100 ms. For simplicity the EPMA results presented in this paper are qualitative due to the complexity of quantitative EPMA data analysis.

Results The layer of sediment obtained from the sample held at 780 °C for one hour is shown in figure 1. It can be seen that the layer is approximately 150 μm thick. Comparison of the grain structure in the bulk of the material with the structure in the sedimented layer indicates that the grain size in the bulk of the material was larger, see figure 1a.

30 μ m

300 μ m

a)

b)

Figure 1. Micrographs showing the layer of sediment produced during stage one of the technique and it’s effectiveness in use with magnesium. a) Low magnification showing presence of sediment layer at the base of the sample. b) High magnification showing the presence of particles within the layer and a refined grain size.

As Zr was removed from the samples, by the sedimentation of particles, the grain size increased steadily. The sample held at 780 °C had a grain size of 222 μm. Following further sedimentation at 760, 740, 720 and 700 °C the grain size of the samples increased to 694, 902, 966, and 1176 μm respectively. The sedimented layers produced by the control experiment are shown in figure 2. It can be seen that following the initial sedimentation at 780 °C a number of particles settled to the base of the sample, resulting in refinement adjacent to the particles. Remelting the sample and holding the liquid metal at 780 °C for a further hour did not result in anyfurther sedimentation of particles, see figure 2b). Remelting the sample by heating to 780 °C and then cooling to 720 °C did cause further particles to form in the melt and settle to the base of the sample as seen in figure 2c. 100 μ m

a)

b)

c)

Figure 2. Base of samples produced during control experiment. a) Sedimentation at 780 °C after initial casting and grain refinement. b) Further sedimentation at 780 °C. c) Further sedimentation at 720 °C.

245 EPMA of the sedimented layer formed in the sample cooled to 720 °C indicated that there was an association of Fe and Si to the Zr present in the sample, see figure 3. The elemental maps indicate that there were high concentrations of Zr, Si and Fe at the grain boundaries of the sedimented layer. There is also evidence of Zr, Si and Fe present at the centre of a few grains. The structures revealed by EPMA can be interpreted as nucleating Zr particles located at the centre of the grains and non nucleating particles that have been pushed to the grain boundaries following solidification. a)

b)

50μm

c)

d)

50μm

50μm

50μm

Figure 3. Elemental maps showing the distribution of Zr, Fe and Si in a sedimented sample. a) Zr content. b) Fe content. c) Si content, d) Back scattered electron image. The rings in d) encircle grain refining particles that can also be seen in the other three images.

To obtain further information concerning the nature of the Zr particles present within the sedimented layers of the samples a detailed EPMA study was performed to compare non-nucleating particles present at grain boundaries with a nucleating particle located at the centre of a grain. The results are shown in figures 4 and 5. EPMA of a nucleating particle is shown in figure 4a-d. The particle was approximately circular in cross-section. The secondary electron image shown in figure 4a) reveals the presence of internal substructure that appears to have a cored formation radiating from the central region. There was a distribution of smaller particles present in the surrounding matrix. The Zr map shown in figure 4b) indicates the particle to consist largely of Zr. Some smaller precipitates are also present on the Zr map identifying them as also containing some Zr. The Fe map in figure 4d) shows only trace amounts of Fe were present in the particle. It is not felt that the detected Fe was the result of interference from the Zr signal as the Fe distribution does not wholly coincide with the presence of Zr. The Si map in figure 4c) shows that some of the smaller particles contained Si. The Si map indicates that there was Si present in the immediate vicinity of the Zr particle and it is possible that the Si shown is a component of the particle. However a number of smaller particles were present that contain Si and these may be the cause of the Si detected in the EPMA map. 5μm

a)

5μm

b)

5μm

c)

5μm

d)

Figure 4. Detail EPMA of a nucleant particle showing a) Secondary electron image, b) Zr distribution, c) Si distribution, d) Fe distribution.

246 EPMA analysis of particles present at the grain boundaries within the sedimented layer of the sample are shown in figure 5. The particles were much smaller, <1μm, and formed agglomerations as they were pushed to the edges of the solidifying grains. The elemental maps show significant levels of Zr, Si and Fe within the particles. 10μm

10μm

10μm

10μm

Figure 5. Detail EPMA of non-nucleant particle showing a) Secondary electron image, b) Zr distribution, c) Si distribution, d) Fe distribution.

The analysis of non-nucleating Zr particles and a Zr nucleant indicated several differences between the two species. The non-nucleating particles were much smaller, <1 μm compared to the nucleant, ≈3 μm. The small particles present at the grain boundaries also appeared to have a higher Si and Fe content. Both the size and the chemical composition differences between the two types of particle may have been influential in determining the viability of the particles as nucleants.

Discussion It has been demonstrated that the sedimentation technique can be used to remove grain refining Zr from the melt. This was demonstrated by the sedimented layer of particles observed in the sample and that the grain size in the layer was finer than that in the bulk of the material. This indicates that the particles present in the layer were active nuclei. It was also shown that as the temperature at which the sedimentation was carried out decreased the grain size in the samples increased. The increase in grain size was the direct result of the removal of Zr in the form of particles from the melt through sedimentation. This result is evidence that the nucleating particles were responsible for grain refinement as a result of the peritectic reaction that occurred as the liquid metal cooled. The control experiment carried out demonstrated that a melt containing no solid state particles large enough to be sedimented in a one hour period was still grain refined by the formation of Zr particles at a lower temperature than the holding temperature, which could still undergo peritectic reaction. This is further evidence that the grain refinement is the result of dissolved Zr forming a dispersion of nuclei throughout the metal which can undergo the peritectic reaction. There were a number of differences between the two species of particles observed during the detailed EPMA study of the sedimented layer. The nucleating particle analysed was larger than the particles found at the grain boundaries, and contained a cored substructure that radiated from the centre. This indicates that the particle underwent growth in the melt. In comparison the particles at the grain boundaries were smaller. The limit of resolution of the microprobe prevented the observation of the structure within these particles. The chemical content of the particles also differed. There was a higher Fe and Si content detected in the grain boundary particles. Only trace amounts of Fe were detected in the nucleating particle and the Si present was mainly associated with finer particles dispersed around the nucleant. It has been stated previously that nucleating Zr particles have been found to contain Fe and Si [6] and that the size of Zr nucleants is of the order of 1μm. It is also known that

247 dissolved Zr has a high affinity for Fe in molten Mg, as Zr additions are used in commercial practice to refine Fe contaminated melts. It could be hypothesised that the Zr particles observed at the grain boundaries in the sedimented samples formed on clusters of Fe and/or Si atoms. Some of these fine Zr particles may have then grown until they became viable nucleation sites for magnesium. This is in agreement with observations made by Tamura [7]. The trace amounts of Fe and Si detected in the nucleating particle may have been the Fe/Si that nucleated the particle. The small Zr particles present at the grain boundaries may not have been viable nuclei due to their small size, i.e. less than the critical radius for nucleation. The high Fe and Si content in the particles may have caused the surface chemistry of the particles to be altered, upsetting the close match between the lattice spacing that usually exists between Zr and Mg. If the above theory is accepted it would begin to further explain the nucleation process of Zr in Mg. It would appear that very fine Zr particles form in the melt as a result of nucleation on Fe and Si present as trace elements in the melt that act as a catalyst for the formation of the precipitates. These particles then grow as Zr is adsorbed on to their surface, until a critical size is reached, at which point they become a viable nucleant for solid Mg. This theory would help to explain how Zr could act as an efficient nucleant in alloys where the concentration of Zr is below the peritectic level of the equilibrium binary phase diagram as stated by [7, 8, 9, 10].

Summary The grain refinement of Mg via the addition of Zr is dependant on the formation of Zr particles in the melt which can undergo a peritectic reaction. This phenomenon is aided by the presence Fe and Si impurities that act as nucleants or a catalyst for the formation of the Zr particles.

Acknowledgements The authors would like to thank Magnesium Elektron for their support and funding for this work, also Mr. Chris Salter of the Materials Science Department, Oxford University for his assistance with the microprobe analysis.

References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]

F. Sauerwald, Z. Anorg. Allg. Chem. 1947, 255. Y. Tamura, J. Japn Inst Light Metals 1998, 48, 4. E.F. Emley, Principles of Magnesium Technology, 1996, Pergamon Press. K.A. O’Reilly, B. Cantor, and P. Enright, Scripta Metallurgica et Materialia 1993, 28. S. Berry, Part II Thesis 1995, Department of Materials, Oxford University. G.A. Fowler, Magnesium Elektron Internal report 1991. Y. Tamura, J. Japan Inst. Light Metals 1998,48. M. Bamberger, Mat. Sci. Tech. 2001, 17. Y. Tamura and T. Motegi, J. Light Metal Weld Constr. 1998, Vol.48. Y.C. Lee, A.K. Dahle, and D.H. StJohn, Met. & Mat. Trans. A 2000, Vol.31A.

Superplasticity in Hot-rolled Magnesium Alloy AZ31 Sheets P. Zhang, J. Wendt, V. Šupik, J. Zhang, F. Holländer Lehrstuhl Metallkunde und Werkstofftechnik, BTU-Cottbus, Cottbus, Germany

1

Introduction

Magnesium alloys, due to their low densities, are particularly attractive for automotive industry for weight saving [1, 2]. Currently, most of magnesium components in automobile are produced by die-casting process, because the high productivity of die-casting process compensates the relatively high cost of magnesium metal. Magnesium alloys AZ91 and AM60 are the primary die-cast alloys for room temperature application, and constitute about 90% of all structural applications of magnesium, because they exhibit superior castability, good corrosion resistance and a good balance between strength and ducitility [3]. In the last a few years, a substantial effort has been made to fabricate magnesium components by mechanical forming, because the components produced in this way have a considerable ductility and specific strength. However, because of the limitation of slip systems in HCP crystal structure, magnesium alloys usually exhibit a poor workability and limited ductility at room temperature. Recently, it is found that magnesium alloys such as AZ91, AZ61 and ZK61 which produced by a primary forming process (rolling, extrusion) exhibit a superplastic behavior due to the grain refinement [4-7]. These results show that it is promising to apply the superplastic forming processing into practical magnesium applications. In present paper, superplastic behavior of magnesium alloy AZ31 sheets produced by hot-rolling was investigated at temperatures of 350-475 °C and strain -5 -1 -2 -1 rates in the range of 8 × 10 s to 2 × 10 s . The change in microstructure of AZ31 sheets during superplastic deformation was also examined.

2

Experimental

The material used in present study was a commercial magnesium alloy AZ31 (Mg-3 wt%Al-1 wt% Zn- 0.3 wt% Mn). The alloy was cast by Otto Fuchs Metallwerke, Meinerzhagen, Germany, and then extruded to a bar with thickness of 20 mm by Honsel AG, Meschede, Germany. The extruded bar was pre-heated at 400 °C for 0.5 h, and finally rolled to a sheet with thickness of 1.2 mm in a rolling mill with hot-rolls of 300 °C at Brandenburgische Technische Universitaet Cottbus, Germany. The total reduction in thickness by rolling is about 94%, reduction per-pass is about 10-20%. Tensile tests were conducted over a range of temperatures from 350 to 475 °C and at initial -5 -1 -2 -1 strain rates from 8.2 × 10 s to 1.6 × 10 s . Specimens for tensile tests with a gauge length of 10 mm and cross-section area of 1.2 × 4 mm2 were machined from the as-rolled sheets without any additional treatments. The tensile axes were parallel to the rolling direction. Light optical microscopy (LOM) was employed for microstructral examination. Grain size was measured by the lineal intercept method. Polarized light was used to reveal orienta-

249 tion contrast. Specimens for LOM were sectioned, cold mounted, ground, polished and then etched in acetic picral (10 ml acetic acid + 100 ml 4% picric acid in ethanol + 5 ml H2O).

3

Experimental Results

The microstructure of hot-rolled magnesium sheets is shown in Fig. 1. One can see that the hot-rolled magnesium sheet has a mixed structure of fine (7 μm) and coarse (20 μm) grains. The fine grains are located in shear bands which inclined to a angle of about 45° with the rolling direction. Volume fraction of the coarse grains was determined to be about 0.7.

Figure 1. Microstructure of hot-rolled AZ31 sheets. The rolling direction is horizontal.

Fig. 2 shows a typical plot of true stress versus elongation for hot-rolled AZ31 specimens deformed at initial strain rate of 1.6 × 10-4 s-1. One can see that, at test temperature of 350 °C, the true flow stress decreases with the increase of elongation. This is possibly due to the dynamic recrystallization (DRX) during deformation. However, such a decrease was not observed for the stress-elongation curves at 450 and 475 °C. On the contrary, a slight increase in flow stress with the increase of elongation was found. This is probably because of a rapid grain growth concurrently with DRX. Fig. 3a shows the elongation-to-failure versus initial strain rate for hot-rolled AZ31. It is found that the specimens deformed at lower rates and higher temperatures exhibit high elongations to failure. A maximum elongation of 350% was achieved at a low strain rate -4 -1 of 1.6×10 s and a high temperature of 475 °C. Fig. 3b shows a series of AZ31 specimens tested up to failure at an initial strain rate of 1.6 × 10-4 s-1 and different temperatures. All the specimens failed by plastic necking. Fig. 4a shows the flow stress at an elongation of 10% against strain rate curves at different temperatures. The mean apparent strain rate sensitive exponent ma obtained in present work is about 0.35–0.42. There is a tendency of slight increase in ma with the increase of temperature. These values of ma are measured at an elongation of 10%, and not affected by grain growth. Fig.4b shows the strain rate against 1/T (T is absolute temperature) plot at σ = 24 MPa. Apparent activation energy can be estimated from the slope of the plot in Fig.4b. The apparent activation energy Qa over 350-475 °C is about 128 kJ/mol, which is close to that for self-diffusion of magnesium 135 kJ/mol [8]. The similar values of ma and Qa were also found in a hot-extruded magnesium alloy AZ61 [7].

250 35

strain rate 1.6×10

true stress (MPa)

30

-4

s

-1

350°C

25 20 15

400°C

10

450°C

5 0

475°C 0

100

200

300

elongation (%)

400

Figure 2. True stress versus elongation over 350-475 °C at an initial strain rate of 1.6 × 10 s . -4

elongation-to-failure (%)

350

-1

350°C 400°C 450°C 475°C

250

150

50

-4

10

(a)

-3

-2

10 10 -1 strain rate (s )

(b)

Figure 3. (a) elongation-to-failure versus initial strain rate over 350-475 °C, (b) fractured specimens deformed to failure at 1.6 × 10-4 s-1 and different temperatures.

Fig. 5 shows optical micrographs of the grip and gauge regions in tensile specimens up to elongations of 60% and 250% (failure) at 400 °C and 8.2 × 10-5 s-1. Compared to the initial microstructure, the static grain growth occurred in the grips (Fig.5a). The small grains in shear bands seem to more easily grow than those coarse grains, however, the mixed structure can be still observed. In gauge region where the dynamic grain growth occurred, it is expected the grain size should be larger than that in grip region. On the contrary, a uniformed and relatively fine grain structure can be seen when the specimen was deformed up to an elongation of 60% (Fig. 5b). This result indicates that the DRX occurred, and resulted in a grain refinement in the gauge region during deformation. The microstructure of a tensile specimen after elongation-to-failure test is shown in Figs. 5c and 5d. No significant change in microstructure in the grip region could be seen with the increase in elongation. In contrast to the grip region, a uniformed coarse grain structure was found in gauge region (Fig. 5d), indicating that a rapidly dynamic grain growth took place in the gauge region. Moreover, no twinning was observed in all deformed specimens.

251 -2

10

stress (MPa)

450°C 475°C

10

1

-4

-3

-3

10

-4

10

3

10

-1

400°C

strain rate (s )

350°C

100

-2

0.0013

10 10 -1 strain rate (s )

(a)

σ =24 MPa

Q a=128 kJ/mol 0.0014

0.0015 -1 1/T (k )

0.0016

(b)

Figure 4. (a) the flow stress at an elongation of 10% as a function of strain rate, (b) the strain rate as a function of 1/T at σ = 24 MPa.

4

Discussion

The present results show that hot-rolled AZ31 exhibits a superplastic behavior at 350475 °C. A maximum elongation of 350% was achieved at 475 °C and a strain rate of 1.6 × 10-4 s-1. It is well accepted that grain boundary sliding (GBS) is the dominant deformation mechanism for superplastic flow [9]. A ma = 0.5 and grain boundary diffusion activation energy are generally associated with GBS. However, the mean apparent strain rate sensitive exponent ma obtained in this study is about 0.35–0.42, which is lower than 0.5. In addition, the apparent activation energy over 350-475 °C is about 128 kJ/mol, which is close to that for self-diffusion of magnesium 135 kJ/mol rather than the grain boundary diffusion activation energy. The values of ma (0.35-0.42) and Qa (128 kJ/mol) suggest that the deformation mechanism in AZ31 is possibly a combination of grain boundary sliding and viscous glide mechanism. Considering the mixed microstructure of fine and coarse grains in AZ31 sheets, one can expect that the fine and coarse grains may deform in a different way, i.e., the fine grains deformed by GBS, whereas coarse grains deformed by viscous glide. Due to the higher volume friction of coarse grains, viscous glide of dislocations by lattice diffusion acts as a rate controlling mechanism, and contributes the most part of deformation. With the aggregation of deformation strain, DRX produced a grain refinement in coarse grains, and made the structure uniformed, as shown in Fig. 5b. Hence, the contribution of GBS to total deformation should increase. However, the grain refinement by DRX is not effective in AZ31 due to a rapid grain growth concurrently with DRX at high temperature. As a result, the contribution of GBS further decreases, and the contribution of viscous glide increases. Furthermore, the grain growth resulted in an increase in flow stress (Fig. 2) and a decrease in ma, and consequently, the specimen with coarse grains fractured by necking (Fig. 5d). It is noted that grain growth can be suppressed at lower temperatures. Accordingly, one can expect a fined grain structure by DRX at lower temperatures. This effect has been in fact observed. A refined grain structure of 6 μm has been obtained in AZ31 at a low temperature of 250 °C by Tan et al. [10], however, the specimen fractured with an elongation of 140%, due to cavitation failure. Tan et al. [10] have proposed a ‘two-stage deformation’ method to improve superplastic ability of AZ31, i.e., firstly, specimen deformed at

252 low temperature to obtain refined grains by DRX, and subsequent deformation was carried out at high temperature to take advantage of refined grains. Nevertheless, the improvement is limited. The maximum elongation at 400 °C can only be enhanced from 250% to 320% by the ‘two-stage deformation’.

(a)

(b)

(c)

(d)

Figure 5. Microstructure in the grip and gauge regions of tensile specimens up to different elongations at 400 °C and 8.2 × 10-5 s-1. (a) grip region at an elongation of 60%, (b) gauge region at an elongation of 60%, (c) grip region at an elongation of 250%, (d) gauge region at an elongation of 250%. Tensile axes are horizontal.

In contrast, an effective way to inhibit grain growth is to introduce precipitates by alloying elements [4-6]. Mohri et al. [5] used a hot-rolled AZ91 sheet, they achieved a maximum elongation of 604% at 1.5 × 10-3 s-1 and 300°C. More recently, Matsubara et al. [11] obtained a maximum elongation of 840% at 200°C and 3.3 × 10-4 s-1 in an ECAPed AZ91 with fine grains of 0.7 μm. Note that superplastic metals are often formed by blow forming, therefore, superplastic metal sheets processed by rolling are needed for superplastic blow forming. Hitherto, however, most of magnesium alloys with a good superplastic ability have been produced by hot extrusion or ECAP, because the high content of alloying elements in these alloys makes the rolling process difficult.

253

5

Conclusions

(1) Superplastic behavior of magnesium alloy AZ31 sheets produced by hot-rolling was investigated at temperatures from 350 to 475 °C and strain rates in the range of 8 × 10-5 s-1 to 2 × 10-2 s-1. A maximum elongation of 350% was achieved at 475 °C and a strain rate of 1.6 × 10-4 s-1. (2) The mean apparent strain rate sensitive exponent ma obtained in this study is about 0.35 - 0.42. The apparent activation energy is about 128 kJ/mol. These values of ma and Qa suggest that the deformation mechanism is possibly a combination of grain boundary sliding (GBS) and viscous glide. The viscous glide of dislocations probably acts as a rate controlling mechanism. (3) Initial microstructure of hot-rolled AZ31 has a mixed structure of fine and coarse grains. During superplastic deformation, dynamic recrystallization (DRX) and concurrent grain growth occurred, consequently, the rapid grain growth resulted in a failure of specimens by necking.

6

Acknowledgements

Thanks are due to the BMBF for financial support of this work.

7

References

[1] Luo, A., Pekguleryuz, M.O., J. Mater. Sci., 1994. 29, 5259-5271. [2] Mordike, B.L., Ebert, T., Mater. Sci. Eng., 2001. A302, 37. [3] Avadesian, M.M., Baker, H., ASM Specialty Handbook: Magnsium and magnesium alloys. Marerials Park, OH: ASM International. 1999. [4] Kim, W.J., Chung, S.W., An, C.W., Higashi, K., J. Mater. Sci. Lett., 2001. 20, 16351637. [5] Mohri, T., Mabuchi, M., Nakamura, M., Asahina, T., Iwasaki, H., Aizawa, T., Higashi, K., Mater. Sci. Eng., 2000. A290, 139-144. [6] Watanabe, H., Mukai, T., Ishikawa, K., Mabuchi, M., Higashi, K., Mater. Sci. Eng., 2001. A307, 119. [7] Wang, Y.N., Huang, J.C., Scripta Mater., 2003. 48, 1117-1122. [8] Frost, H.J., Ashby, M.F., Deformation mechanism map. Oxford: Pergamon Press, 1982. [9] Valiev, R.Z., Langdon, T. G., Acta metall. mater., 1993. 41, 949. [10] Tan J.C. Tan, M.J., Scripta Mater., 2002. 47, 101-106. [11] Matsubara, K.M., Y. Horita, Z. Langdon, T.G., Acta mater., 2003. 51, 3073-3084.

Extrusion of Magnesium Profiles and Microstructure Characterization Sören Müller*, Ingwer Denks*, Klaus Müller* and Anke Pyzalla** * Technical University Berlin, Berlin, Germany ** Vienna Technical University, Vienna, Austria

1

Abstract

In order to employ extruded profiles of the magnesium alloys AZ31 which are a considerable alternative to aluminium in light-weight constructions, the extrusion process itself as well as the mechanical properties of the profiles have to be optimized. The extrusion trials on the Mg alloy AZ31 were carried out on the 8 MN horizontal extrusion press at the Extrusion R & D Center (ERDC) at the Institute for Material Science and Technology of the Technical University Berlin, Germany. Products with different shapes cross-sections, but with similar extrusion ratios were extruded. The billet temperature, extrusion speed and cooling rates were variied. The mechanical properties of the extrudates such as compression/tensile strength, yield strength and fracture elongation were determined. Also texture analysis, microhardness measurements, SEM/ TEM and optical microscopy for grain size measurement were carried out. An influence of grain size as well as texture and microstructure on the mechanical properties could be detected. Analyses of different areas of the discard (billet, deformation zone, product) showed the development of the microstructure and the texture. The influence of the product shape on the extrusion process and the mechanical properties was analysed by a comparison of extrudates produced with the different extrusion forces which have significantly different microstructures. Further, the correlation between the precipitation state and the texture development was determined.

2

Introduction

Although magnesium components are increasingly employed in the aerospace and automobile industry these components are almost exclusively casted parts [1]. Because of the hexagonal (hdp) lattice of the magnesium its deformation mechanisms are different to those of other light-weight metals such as aluminium (fcc). The difference in deformation mechanisms has an influence on the extrusion process itself as well as on the resulting product. In the beginning of the extrusion of magnesium wrought alloys a high peak force can be observed before the process stabilizes and the forces reach a stationary condition [2]. The extruded products show a mechanical anisotropy which has an important impact on the application of magnesium profiles because the yield point under compressive load is significantly lower than under tensile load [3,4]. Therefore it is necessary to obtain a clear understanding of the relationships between the extrusion process (and its parameters), the microstructure and the mechanical properties.

255

3

Experimental procedure

Magnesium billets produced by continuous casting were extruded indirectly at the R&D Center for Extrusion of the Technical University Berlin. The extrusion trials were carried out at a temperature of 300 °C (billet and container) with various speeds from 1 m/min up to 30 m/min and air respectively water cooling. Since the extrusion press is equipped with load cells the occurring extrusion forces(total force, die force) can be recorded. Thereby comparisons of the necessary extrusion forces in the different extrusion trials were possible. Afterwards samples were taken out of the products as well as from determined positions within the discards. Optical microscopy was performed on specimens from the billet, the discard and the product in order to receive an overview over the change in the microstructure during the deformation process. In addition to optical microscopy SEM and TEM were used for analysing the deformed microstructure, precipitation state and segregations in the discard and the billets. Synchrotron radiation was used for profile analyses in specific regions of the discard. The results were compared to those obtained by optical microscopy and TEM. In order to determine the influence of the extrusion parameters (extrusion speed and cooling conditions) on the mechanical properties of the products X-ray texture analyses and mechanical testing were performed.

4

Results

4.1

Billet

The AZ31 billets showed a typical cast structure with grain sizes up to several millimeters. Figure 1 shows the microstructure in the billet. It contains dendrites and a surrounding interdendritic area. The dendrites are seperated by grain boundaries that are running through the interdendritic area (Fig. 2).

Figure 1. Cast material containing dendrites

Figure 2. Interdendritic area with grain boundaries

Precipitates have only been observed in the indendritic areas. Within this area the precipitations are distributed homogeneously, a concentration of precipitates along the grain boundaries was not observed.

256 The cast material contains two different types of precipitates. One type of precipitates is γ (Mg17Al12). The other type contains Mn in addition to Mg and Al. Twins can also be found in the as-cast material. These twins are parallel within each grain and they run through the interdendritic area as well as through the dendrite. 4.2

Discard

In the beginning of the extrusion process the billet is compressed in the container before it is extruded through the die. These different stages of deformation can still be seen in the discard after the process (Figure 3). 110mm

Zone of the beginning multi-axial deformation

Compression zone

Deformation zone

Figure 3. Discard

Figure 4 shows that twinning is the predominant deformation mechanism in the compression zone. Very large twins that cross through whole grains as well as small twins within the space between the large twins are distinguishable. In the area where stronger and multi-axial deformation occurs the boundaries of the twins are blurred. The blurring of the twin boundaries indicates a dislocation movement inside and ouside of the twins (Figure 5). New twins are only very rarely formed in this area. The appearance of small shear bands which are surrounded by recrystallized areas can be observed.

200µm Figure 4. Twinning in the compression zone

50µm Figure 5. Blurred twins in zones with a higher degree of deformation

257 In the deformation zone of the discard where maximum deformation occurs the number and the size of the shear bands increases (Figure 6) The shear bands cross the interdendritic and the dendritic area as well as the grain boundaries. Closer to the die the appearance and the dimension of the shear bands increases. Far away from the die the shear bands are not wider than 50 μm whereas close to the die shear bands with a width of 200 μm were observed. With increasing deformation the width of the recrystallized areas along the shear bands also increases (Figure 7). Recrystallization starts at the grain boundaries. Toward the die the grains are elongated so that recrystallized areas can also be seen between the shear bands.

20µm

100µm Figure 6. Shear bands

4.3

Figure 7. Recrystallization along the shear bands

Product

The microstructure of the product can mainly be divided into three different parts (Figure 8). Area I is a recrystallized area which consists of grains with a grain size between 5 μm – 20 μm. The precipitates between the grains indicate this area as the interdendritic area. In contrast to the cast material in the product only one type of precipitates was observed. These are the Mn-rich precipitations. The Mg17Al12 precipitations have been dissolved. Area II is a dendritic area that is also recrystallized but with an average grain size of 5 μm. Area III contains regions of various dimensions that are not recrystallized or only partly recrystallized. In some of these regions significant twinning can be observed. These twins are mainly oriented at 45° to the extrusion direction (Figure 9). Area II Area III

Area I

10µm Figure 8. Microstructure of the extrudates

10µm Figure 9. Twinning in the extrudate

258

5

Process modification

3,5

7

3

6

2,5

5

2

4

1,5

3

Total Force

1

2

Die Force

0,5

Ram spedd [mm/s]

Force [MN]

After the characterization of the microstructure and the deformation mechanisms in the discard and the products as described in paragraphs 2 - 4.3 new extrusion trials were performed. Because the deformation of small recrystallized and statiscally distributed grains is easier than the deformation of large grains additional extrusion trials with different process conditions were performed. In one set of extrusion trials the billet temperature was increased up to 380 °C. In a second set of extrusion trials a speed profile with increasing speed throughout the process was used. The extrusion diagrams of these trials (Figure 10 and Figure 11) show that due to the increase in billet temperature as well as due to the speed profile a significant reduction of the peak force in the beginning could be achieved.

1

Ram Speed

0

0 0

50

100

150

200

250

300

Ram displacement [mm]

3,5

14

3

12

2,5

10

2

8

1,5

Total Force

1

Die Force

0,5

Ram Speed

0

6 4 2 0

0

50

100

150

200

Ram displacement [mm] Figure 11. Extrusion diagram with a speed profile, billet temperature 300 °C

250

300

Ram speed [mm/s]

Force [MN]

Figure 10. Extrusion diagram with a billet temperature of 380 °C

259

6

Conclusion and Outlook

The deformation mechanisms of the magnesium alloy AZ31 were investigated at the different stages of an indirect extrusion process. In the beginning of the extrusion when the billet is compressed in the container twinning is the predominant deformation mechanism. As deformation increases during the actual extrusion shear bands are formed that are surrounded by recrystallized grains. With the knowledge of the results from the different analyses a specific variation of the extrusion parameters could be deducted. New extrusion trials were carried out where a reduction of the peak force in the beginning of the extrusion process could be achieved. The analysis of the influence of the modified extrusion process on the microstructure and mechanical properties of the product is the aim of ongoing research.

7

Acknowledgements

This work is supported by Otto Fuchs Metallwerke, Meinerzhagen, Germany and the DFG (DFG-Project Py 9/6-1).

8

References

[1] D. Brungs, H. Gers, A. Mertz, „Recent Advances in Magnesium Processing Technologies“, Thermec 2000, Las Vegas , NV, p. 52. [2] S. Müller, A. Pyzalla, Strukturuntersuchungen an indirekt stranggepressten Magnesiumlegierung des Typs AZ, in: Tagungsband 9. Sächsische Fachtagung Umformtechnik, p. 243-253, Dresden, 2002. [3] J. Becker, G. Fischer, Strangpressen und Eigenschaften von Profilen aus hochfesten Magnesium-Legierungen, G. Fischer (Hrsg.), Strangpressen, DGM Informationsgesellschaft GmbH, Frankfurt a. M., 1997. [4] K. Müller, „Direct and Indirect Extrusion of AZ 31“, Magnesium Technology 2002, H.I. Kaplan (Ed.), p. 187-192, Seattle, USA, 2002.

Investigating the Plastic Deformation Behaviour of Magnesium Alloy AZ31 by Plane Strain Forging of U and H Sections M.S.Yong, S.C.V.Lim Singapore Institute of Manufacturing Technology, Singapore

1

Introduction

Research in magnesium has received greater attention since mid 90's, mainly due to the increased use of the material for automobile. Significant achievement has been made in the areas of machinery, alloy development and applications in automotive industry. Nonetheless, the forming technologies for magnesium alloy, particularly bulk forging is not widely used because of limited knowledge about the technology and understanding of the process parameters [1]. Although there has been a series of forged magnesium components developed in the recent years, very little detailed work were reported so far in those areas. A number of studies were conducted on the plastic deformation behaviour of axisymmetric forging of magnesium alloy [2] [3]. However, many products are far from axisymmetric in shape. Therefore, this paper will report on the work conducted to evaluate the plastic deformation of magnesium strip using the plane strain forging method. Finite element analysis and microstructure were used to correlate the plasticity flow of the material under high load and temperature.

2

Experimental Procedure

A 2D plane-strain model consists of a forming punch, billet material and die insert were used for the finite element simulation. The different forging loads and temperatures experienced by AZ31 billet material was modelled as a strain-hardening plastic body governed by the power law σ = Kεn. The material constants K and n at different forming temperatures of the material were obtained from internal reports of AZ31 and are shown in Table 1. ANSYS/LS-DYNA simulation software was used for the processing of the model. The coefficient of friction value (μ) for the experiment was found to be 0.12 [4]. The value is subsequently used in the simulation for the generation of load displacement graphs found in this paper. Table 1. Material Constants of AZ31

Temperature (°C)

150 250 300

σ = Kεn K

n

178 85 59

0.1 0.024 0.021

261 The alloy used in this investigation was a commercial grade of magnesium alloy AZ31 of composition presented in Table 2. Flat blanks of 100 mm (length) by 25 mm (width) by 5 mm (thickness) were used in the plane strain forging experiment. The blanks were deformed to U and H sections of 1mm thickness, and 2 mm flange thickness, as shown in Figure 1. The experimental setup can be seen in Figure 2. The processing parameters considered for warm forging were: forging load of 56 ton, forging temperature of 150 °C, 250 °C and 300 °C. The blanks were coated with graphite to act as a layer of dry lubricant. Immediately after forging the forged samples were water quenched to arrest any possible changes in microstructure. Table 2. The specification of magnesium alloy AZ31 used in the investigation

Material

Al

Zn

Mn

Si

AZ31

2.8%

0.9 %

0.2%

<0.01% <0.01% 0.004% Balance

56 ton

0

10 20 mm

53 ton

50 ton

o

250 C

40 ton

Cu

Fe

Mg

(a)

30 ton (b)

o

250 C

o

300 C

0

10 20 mm

Figure 1. Photograph of the formed U and H section using the (a) backward only, and (b) forward & backward extrusion process respectively.

3

Results and Discussion

3.1

Load-Displacement Curves

Figure 2. Photograph of the tooling used for warm forging AZ31 alloy.

Using μ = 0.12, simulations of forging at 150 °C, 250 °C and 300 °C were carried out and the forging load vs punch displacement are shown in Figure 3 along with the empirical results for 250 °C and 300 °C. Working within the constraint (maximum capacity) of the hydraulic press, a load of 56 ton was used. When heated at 150 °C the blank displayed minimum plastic deformation compared to the blank forged at 250 °C and 300 °C under the same load of 56 ton. Generally, the load-displacement curves of the experimental results showed good agreement to that predicted by the FEM simulation. It was also found that the 56 ton loading could only deform thickness from 5 mm to 1.9 mm (i.e. punch displacement = 3.1 mm) at forging temperatures of 250 °C. By increasing the forging temperature to 300 °C, the billet could be deformed to 1mm in thickness with a mere 50 ton loading. The requirement of a much higher load needed to forge the billet at the temperature of 150 °C as compared to that above 250 °C, can be attributed to the HCP structured magnesium, where the deformation of magnesium above 225 °C causes additional slip planes { 10 11 } to become operative [5].

262 180

200

160 120 100

o

250 C FEA

80

o Expt. 250 C

60

o

300 C FEA

40

140 120

FEA 300oC F&BW process

o

FEA 250 C BW process

100 80 60 40

o

Expt. 300 C

20

o

160 Load (tons)

Load (Tons)

140

FEA 250 C F&BW process

180

150oC FEA

o

FEA 300 C BW process

20

0

0

0

1

2

3

4

5

Punch Displacement (mm)

Figure 3. Graph showing the load vs punch displacement for the simulations and experiments at different forging temperatures for the backward extrusion process.

0

1

2

3

4

5

Punch Displacement (mm)

Figure 4. Graph showing the load vs punch displacement for the simulations at different forging temperatures for the forward and backward extrusion processes.

Forging load vs punch displacement curves obtained at simulations for forging conditions at 250 °C and 300 °C for the close-die forward and backward extrusion process (Figure 4) were compared to that of the open-die backward extrusion process (Figure 3). The profile of the curves is quite similar in the two processes, except the large build-up of load at the final stage for the forward and backward extrusion process. This is attributed to material filling up the cavity as this process is a close die process. It is observed that a smaller load is required to forge the part from 5 mm thickness to 2 mm (punch displacement = 3 mm) for the respective different temperature used as compared to the parts forged in the backward extrusion process (Figure 4). The smaller load required was attributed to the greater ease for material to flow as the material has additional avenue to move. It can be seen that approximately the same amount load was required for the punch displacement of 3 to 3.5 mm regardless of which process used. Thus, the benefit of the extra cavity in the forward and backward extrusion process was evident only for the punch displacements prior to 3 mm. 3.2

Microstructure of U specimen forged at 150 °C with a load of 56 ton

Figure 5 compares the macrograph of the part forged at 150 °C with a load of 56 ton with the strain distribution profile obtained through FEA. It was observed that the microstructure pattern coincides accurately with the strain distribution. At area A, where the strain was relatively the highest, we do observe a high concentration of deformation twins/ twinning lines (Figure 6a). As the magnitude of strain was low, the twins were generally long, thin and lamellae. In area B, twinning lines were also observed but only along the shear band, indicating that the strain was higher in the shear band region as compared to the region around it (Figure 6b). Such observation of twinning lines in the microstructure is due to the limited slip systems of the magnesium HCP structure at temperature below 225 °C [5]. In areas experiencing low strain, the microstructure generally remains undeformed (Figure 7).

263 Area A

Area A

Area B

Figure 5. Macrograph of part forged at 150 °C matched with the strain distribution diagram as predicted by FEA.

(a)

(b)

50 µm

Figure 6. Micrographs showing the microstructures with deformation twins in (a) Area A, (b) Area B.

3.3

100 µm

50 µm

Figure 7. Micrograph showing the microstructure without deformation twins at areas of low strain.

Microstructure of U specimen forged at 250 °C with a load of 56 ton

As the forging temperature increased to 250 °C, the billet could be forged much easier than at 150 °C. Besides improved forgability, the microstructural evolution was observed to be different. Figure 8 shows a good correlation of the macrograph the U specimen forged at 250 °C using load of 56 ton with the strain distribution diagram as calculated by FEA.

50 µm

ε = 1.42 – 1.84 ε = 1.98 – 2.26

50 µm

ε = 1.0 – 1.28

ε = 2.26 – 3.09 Figure 8. Macrograph of part forged at 250 °C using load of 56 ton matched with strain distribution diagram obtained from FEA showing the non-homogeneous strain distribution.

264 It was observed that there were no deformation twins/twinning lines in the structure. Generally, fine equiaxed grains were observed and the strain values were in the region of 1.5 and above (FEA strain distribution diagrams). The fine equiaxed grains were developed through dynamic recrystallization during the forging process. In regions of relatively lower strain i.e. values of 1.0 to 1.4, it was observed that certain areas have elongated grains surrounded by fine equiaxed grains. As the strain distribution was non-uniform, it was found that the microstructure was non-homogeneous throughout the forged part as well. This can be attributed to dynamic recrystallization being strain dependent [6] which would influence the grain structure formed during the forging process. Generally, by forging the billet at 250 °C using a load of 56 ton, a part with a more refined microstructure was obtained as compared to the initial undeformed stage. The strain diagram obtained from the FEA can act as a good indicator of the grain size with region of high strain generally having smaller refined grains. 3.4

Microstructure of H specimen forged at 150 °C & 250 °C

Generally, the microstructures of the parts forged at 150 °C and 250 °C under the forward and backward extrusion process were similar to those using the backward extrusion process. For the part forged at 150 °C, deformation twins were observed in areas of considerable strain (ε = 0.1 and above) as shown in Figure 9. The concentration of twins increased in regions of higher strain. In areas of low strain (ε = 0.05 and below), grains identical to those found in the initial undeformed structure were observed. For the parts forged at 250 °C, the microstructures were fine, recrystallized grains with no twining lines (Figure 10). The size of the grains and their structure were very dependent on the strain distribution too. The micrographs revealed that the forward and backward extrusion process did not have a great influence on the microstructure of the part formed as compared to the temperature of forging. (a)

(b)

50 μm

50 μm

Figure 9. Representative micrographs showing the microstructure at (a) area of relatively higher strain value as compared to (b) area with low strain value for the part forged at 150 °C using the forward and backward extrusion process.

50 μm

Figure 10. Representative micrographs showing the microstructure of fine recrystallized grains for the part forged at 250 °C using the forward and backward extrusion process.

265

4

Conclusions

1. A smaller load was required to forge the part from 5 mm thickness to 2 mm (punch displacement up to 3mm). The smaller load required could be attributed to the greater ease for material to flow as the material had additional avenue to move. 2. The microstructure of the U specimen forged at 150 °C under the “backward” extrusion process showed a high concentration of deformation twins/ twinning lines at region where strain was high. In areas of low strain, the microstructure generally remained the same as in the initial undeformed stage. Twinning lines in the microstructure are due to the limited slip systems of the magnesium HCP structure at temperature below 225 oC for magnesium alloys. 3. No deformation twins/twinning lines were detected in the microstructure of the U specimen forged at a higher temperature of 250 °C under the “backward” extrusion process. Fine equiaxed grains were observed in regions of high strain. As a whole, the formed structures shown a more refined microstructure compared to the initial structure. The strain distribution diagram as generated from the FEA correlated well with the distribution of the grain size in the forged part. Region of high strain generally had smaller refined grains. 4. The microstructure of the H specimens forged under the “forward and backward” extrusion process was very much similar to what was observed for the parts forged using the “backward” extrusion process at 150 °C and 250 °C. In both process, deformation twins were observed at 150 °C, whilst fine and recrystallized grains with no twining lines were observed for the part forged at 250 °C and above.

References [1] Doege, E., Droder, K., Hamm, F.P., Sheet Metal Forming of Magnesium Alloys, Proceedings of IMA-Conference on Magnesium Metallurgy, October 1996, Clermont-Ferrand, France. [2] Chan, C.F, Yong, M.S, Tay, C.J., Shang, H.M., The Influence of Process Parameters on Forged Magnesium Alloys”, International Conference on Advanced Material Processing (iCAMP 2002), 2-4 Dec 2002, Singapore. [3] Kainer, K.U., Doege, E., Janssen, S., Ebert, T., Microstructure, Mechanical Properties and Deformation Behaviour of Extruded Magnesium Alloys”, International Conference on Magnesium alloys and their Applications, DGM' 2000, Pub. Wiley-Vch, Germany, 596-601. [4] Lim, S.C.V., Yong, M.S, Choy C.M., Evaluating Plain-strain Forging of Magnesium Alloy AZ31 using Finite Element Analysis, IC-SEC 2002 International Conference on Scientific and Engineering Computation, 3-5 Dec 2002, Singapore. [5] Emley, E.F., Principles of Magnesium Technology, Pergamon Press, Oxford, New York, 1966, 483-488. [6] Humphrey, F.J. and Hatherly, M., Recrystallization and Related Annealing Phenomena, Pergamon Press, 1996, 177-178.

Application of Isothermal Rolling for Improvement of Sheet Formability of a Magnesium Alloy A. Galiyev, R. Kaibyshev, D. Voronin Institute for Metals Superplasticity Problems, Ufa

1

Introduction

Rolling is the one of promising processes for magnesium technology applications. In comparison to die-castings magnesium sheet metal offers the combination of advantageous properties like good deformability, high strength and ductility. To date, however, very limited data are available for rolling processed magnesium-based materials. It is known that the coarse grain structure of the chill cast material can be transferred into a finer recrystallized microstructure by rolling at temperatures between 480 °C and 300 °C [1]. These rolled material indicate a satisfactory sheet quality and material characteristics. For a subsequent treatment a sheet metal is to produce with good technological properties. The significant improvement of formability can be attained in semi-finished material with very small grain size [2] because the reduction of the grain size results in achieving superplasticity. However, the potential of grain refinement for further improvements of magnesium sheet metal properties is not fully utilized. In the present paper, thermomechanical processing (TMP) was conducted for ZK60 Mg alloy in order to produce fine-grained microstructure and enhance material’s ability to hot rolling. Isothermal rolling was performed on the grain-refined material and the mechanical properties of the rolled Mg alloy were investigated at room and elevated temperatures to evaluate superplastic ability and forming potential of the magnesium sheet metal.

2

Experimental

A ZK60 (Mg-5.8%Zn-0.65%Zr) alloy ingot was subjected to TMP including hot pressing at temperatures from 390 to 330 °C with a total strain of 0.84 followed by isothermal compression at 350 °C with a strain of 1.4. This TMP resulted in the formation of partially recrystallized structure with a grain size in the range of 4-16 μm. Blocks 10 mm in thickness and 45 mm in breadth were cut from the thermomechanically processed ingot with the rolling direction (RD)-transverse direction (TD) plane parallel to the compression axis and isothermal rolled at a temperature of 300 °C with a 70.8% reduction. Microstructures of the thermomechanically processed and rolled specimens were observed by an optical microscopy. To evaluate tensile properties of the alloy, tensile specimens 6 mm in guage length, 2 mm in guage thickness and 3 mm in guage width were machined with the tensile axes perpendicular to the upset direction and parallel to the rolling direction for the thermomechanically processed and rolled specimens, respectively. Tensile tests were carried -3 -1 out in air at room temperature at an initial strain rate 2.8 × 10 s and elevated temperatures in the range from 250 to 300 °C with initial strain rates from 1.4×10-4 to 5.6 × 10-2 s-1.

267

3

Results and Discussion

3.1

Microstructure

Microstructures of the Mg alloy after TMP are shown in Fig. 1. It can be seen that recrystallization occurred by hot deformation processing because the fine, equiaxed grains were partially observed for the specimens after hot pressing and isothermal compression. Systematic observations in the middle and outer parts of the compressed bars revealed more uniform microstructures for the latter regions where significant grain refinement was attained (Table 1). The fine-grained microstructures were inhomogeneous in grain size and scattered from 4 to 16 μm that could be possible due to grain growth during TMP.

Figure 1. Microstructures of the specimens after TMP: (a) in the middle and (b) outer parts of the compressed bar

Figure 2. Microstructures of the specimens after TMP and isothermal rolling: (a) the specimen with coarsegrained microstructure after TMP and (b) the specimen with fine-grained microstructure after TMP. The RD-TD plane is observed

268 Table 1. The recrystallized volume of the Mg alloy for the different parts of the compressed bars

Processing condition hot pressing + isothermal compression isothermal rolling

The middle part of the compressed bars 30%

The outer part of the compressed bars 75%

60%

95%

Figure 2 shows the microstructures of the specimens after TMP and isothermal rolling, where the RD-TD plane is observed. The rolling was carried out for two states of Mg alloy processed by TMP: one with fine-grained microstructure of the specimens taken from the outer part of compressed bar and the other with coarse-grained microstructure of the specimens taken from the middle part. At rolling temperature 300 °C the specimens even with coarse-grained microstructure could be rolled to a thickness of 0.35 mm in without cracking. In this case, the rolling reduction was 96.3%. In the case investigated, the specimens were rolled to reduction of 70.8%. In the all specimens rolled the coarsegrained microstructures remained after TMP were refined further (Table 1) and the microstructures appeared more homogeneous in the isothermal rolled conditions. For the rolled fine-grained Mg alloy the grain refinement was attained over almost a whole surface of the specimen, while for the rolled coarse-grained specimens the fraction of fine grains increased from 30 to 60%. On the other hand, the small grain sizes of the rolled coarse-grained Mg alloy were almost the same as those of the specimens after TMP, while the grain sizes of the rolled fine-grained Mg alloy decreased to the range of 2-11 μm. 3.2

Tensile properties at room temperature

Tensile tests at room temperature were carried out for isothermal compressed bar and isothermal rolled specimens. The results are summarized in Table 2. It can be seen that the compressed Mg alloy with fine-grained microstructure showed higher 0.2% proof stress and Table 2. Tensile properties at room temperature for isothermal compressed bar and isothermal rolled specimens

Material condition

0.2% Proof stress [MPa]

Ultimate tensile strength [MPa]

Elongation to failure [%]

the middle part of the compressed bar the outer part of the compressed bar

176

266

12

193

284

31

rolled coarsegrained specimens rolled finegrained specimens

216

295

31

233

304

34

269 higher ultimate tensile strength than the one with coarse-grained microstructure. Another more important result is that the elongation to failure of the former was much higher than that of the latter. This trend is the same as that shown by Chapman and Wilson [3] for Mg. The rolled specimens were characterized by higher ultimate tensile strength, 0.2% proof stress and elongation to failure compared with those of the specimens prior to rolling. The good mechanical properties are likely attributed to further grain refinement caused by DRX during rolling as evident from the comparison of the data in Tables 1 and 2. As a result the maximum ultimate tensile strength of 304 MPa, 0.2% proof stress of 233 MPa and elongation to failure of 34% were obtained in the rolled fine-grained specimens. 3.3

Superplastic behavior

Figure 3 shows the typical true stress-true strain curves for isothermal compressed bar and isothermal rolled specimens at 250, 275 and 300 °C at an initial strain rate of 5.6 × 10-4 s-1. Only the curves for these conditions are shown since optimum superplastic elongations were found near these temperatures and strain rate. It is apparent from Fig. 3a that the compressed specimens with coarse-grained microstructure exhibited much higher stresses than the specimens with fine-grained microstructure. Isothermal rolling of the coarsegrained specimens resulted in essential decrease in stress (Fig. 3b). This could be attributed to grain refinement during rolling. Essential shift of the peak stress to a higher strain was observed in the rolled specimens. Figure 4 shows the variation of strain rate sensitivity, m, with initial strain rate. For the compressed coarse-grained specimens the m values were low (<0.3) and increased essentially after rolling (0.44 at 5.6 × 10-4 s-1 and 250 °C) (Fig. 4a). It was apparent that superplastic behavior, at which m > 0.3, was achieved in the rolled specimens. For the finegrained specimens the high m values were attained both in compressed and rolled specimens (Fig. 4b). The higher m values of about 0.55 were observed in the rolled specimens at 250 and 275 °C at 5.6 × 10-4 and 1.4 × 10-3 s-1. It should be noted that the further raising of the temperature to 300°C or strain rates to 5.6 × 10-2 s-1 resulted the m values in decrease. (a)

Coarse-grained specimens Fine-grained specimens

40

30

250°C 275°C

True Stress [MPa]

True Stress [MPa]

50

30 250°C 20

275°C

10 0 0.0

300°C 0.5

1.0

1.5

True Strain

25 20 15 10

Coarse-grained specimens Fine-grained specimens

250°C 275°C

250°C 275°C

300°C

5 0 0.0

2.0

(b)

0.5

1.0

1.5

2.0

2.5

True Strain

Figure 3. True stress-true strain curves at 5.6 × 10 s : (a) the specimens after TMP and (b) the specimens after TMP and isothermal rolling -4

-1

270 0.6

(a)

0.4 0.3 0.2 0.1

TMP 250°C 275°C

TMP+Isothermal Rolling 250°C 275°C

0.0

Strain Rate Sensitivity

Strain Rate Sensitivity

0.5

(b)

0.5 0.4 0.3 0.2

TMP 250°C 275°C 300°C

0.1 0.0

-3

10

10 -1

-4

10

TMP+ Isothermal Rolling 250°C 275°C -3

10

300°C

-2

10

-1

-1

Strain rate [s ]

Strain Rate [s ]

Figure 4. Variation of the strain rate sensitivity as a function of strain rate: (a) the specimens with coarsegrained microstructure after TMP and (b) the specimens with fine-grained microstructure after TMP

Variation of elongation to failure as a function of temperature at initial strain rate of 5.6 × 10-4 s-1 is shown in Fig. 5. In general, variation of elongation was found to be corresponding to variation of strain rate sensitivity, m. It can be seen in Fig. 5 that the compressed coarse-grained specimens with low m values showed low elongation of 75 and 82% at 250 and 275 °C, respectively. The isothermal rolled coarse-grained specimens showed essential increase both in m value and elongation. Large elongation of 290 and 340% was attained in the rolled specimens with coarse-grained microstructure at 250 and 275 °C, respectively. This comparison demonstrates a significant improvement of the ductility of coarse-grained Mg alloy isothermally rolled at 300 °C. The compressed fine-grained specimens with high m values showed much larger elongations than the compressed specimens with coarse-grained microstructure. Large elongation of 460% was obtained in the compressed fine-grained specimens at 250 °C. The optimal superplastic behavior was found at 275 °C where maximum elongation of 780% was TMP TMP+Isothermal Rolling

Elongation to Failure [%]

1200 1000

coarse-grained specimens fine-grained specimens

800 600 400 200 0

260

280

300

Temperature [°C] Figure 5. Variation of the elongation to failure as a function of temperature at strain rate of 5.6 × 10 s -4

-1

271 attained. By isothermal rolling still higher superplastic properties of the fine-grained specimens were obtained. At 275 °C elongation for the rolled specimens increased to 920%. At 250 °C an elongation of 1170% was achieved. Essential increase in elongation in the rolled fine-grained specimens was attributed to additional grain refinement occurring during isothermal rolling. The elongation decreased with increasing temperature and it was probably due to grain growth occurring in the fine-grained specimens during testing. By comparison with the specimens before rolling it was well demonstrated that application of isothermal rolling was very important for attainment of significant superplasticity and, apparently, good formability of sheet Mg alloy. Further investigations are required for evaluation of high-strain rate superplasticity and forming potential of the magnesium sheet metal.

5

Conclusions

(1) Thermomechanically processed specimens of ZK60 Mg alloy could be isothermally rolled to reduction of 96.3% in without cracking at 300 °C. (2) The microstructure of rolled specimens was depended on the one after TMP. For thermomechanically fabricated specimens with relatively coarse- and fine-grained microstructures, the volume fraction of fine recrystallized grains increased from 30 to 75 and from 60 to 95%, respectively, after isothermal rolling at 300 °C with a reduction of 70.8%. The size of recrystallized grains was obtained in the range of 4-16 μm for all the thermomechanically processed specimens and the rolled coarse-grained specimens. It was reduced to 2-11 μm in the rolled fine-grained specimens. (3) The strength and elongation to failure of the rolled specimens were larger than those of the non-rolled specimens at room temperature. The high strength and elongation of over 30% were attributed to grain refinement. (4) The specimens with coarse-grained microstructure showed superplastic behavior after isothermal rolling, namely, high elongation of 340% and strain rate sensitivity of 0.36 -4 -1 at 275 °C and initial strain rate of 5.6 × 10 s . (5) Grain refinement in the thermomechanically processed specimens significantly improved superplastic properties of sheet Mg alloy. In the grain-refined ZK60 Mg alloy -4 -1 superplastic elongation of 780% was obtained at initial strain rate of 5.6 × 10 s and a testing temperature of 275 °C. By following isothermal rolling, superplastic elongation of 1170% was attained at lower temperature of 250 °C and the same strain rate. Apparently, good formability should be achieved in a sheet Mg alloy.

6

References

[1] Enss, J.; Everetz, T.; Reier, T.; Juchmann, P., in Proceedings of the 5th International Conference Magnesium Alloys and Their Applications, 27-28 September, 2000, Wiley-VCH 2000, p. 590-595. [2] Mohri, T.; Mabuchi, M.; Nakamura, M.; Asahina, T.; Iwasaki, H.; Aizawa, T.; Higashi, K., Mater. Sci. Eng. 2000, A290, 139-144. [3] Chapman, J. A.; Wilson, D. V., J. Inst. Metals 1962-63, 91, 39-40.

Influence of Al Content (<3%) on the Extrusion Limit Diagram Matthew R. Barnett1, Ji-Yong Yao2, Chris Davies3 CRC for Cast Metals Manufacturing (CAST) 1 Deakin University, Geelong, VIC, Australia 2 University of Queensland, Brisbane, QLD, Australia 3 Monash University, Clayton, VIC, Australia

Introduction One of the impediments to the uptake of wrought magnesium alloys has been their perceived poor forming characteristics when compared to aluminium. Extrusion speeds for magnesium are typically quoted as between 2/3 and 1/3 that of aluminium [1], with a concomitant reduction in productivity. However, it is also apparent that the extrusion of magnesium alloys and the alloys themselves have not been optimised. With CAFE (Corporate Average Fuel Economy) legislation in the US and EU for fuel economy there is an opportunity for the weight reduction of motor vehicle components which takes advantage of the optimised processing of wrought magnesium. This paper describes a methodology which allows us to evaluate wrought alloys and their processing based on laboratory-scale testing coupled with simulation and modelling.

Extrusion Limit Diagrams An effective way of comparing the extrudability of different alloys is through the generation of extrusion limit diagrams [2]. A schematic example of such a diagram in terms of extrusion ram speed and billet temperature is shown in Figure 1. Two limit lines form an envelope enclosing a region where extrusion is possible. At low temperatures, the press capacity imposes a limit on the speed that can be reached. At higher temperatures the incipient melting of an alloy phase causes surface cracking. The maximum ram speed shown is 100 mms-1. Speeds higher than this are unlikely to be encountered on typical extrusion presses due the limitation of the press hydraulics. To aid in the generation of an extrusion limit diagram it is desirable to develop a mathematical framework that enables the rapid estimation of loads and maximum surface temperatures. A simple approach is described below, based on engineering equations readily available in the literature [e.g. 2]. The extrusion pressure is given by: ⎛ ⎞ p = σ ⎜ 0.17 + 1.86 ln R + 4mlb ⎟ d 3 b ⎝ ⎠

(1)

where σ is the mean von Mises effective flow stress, R is the extrusion reduction ratio, m is a friction factor (set to unity here), lb is the billet length and db is the billet diameter.

273

Ram speed (mms-1)

100 Press load limit exceeded

10

1 200

Surface cracking Safe extrusion

300

400

500

600

Billet temperature (°C)

Figure 1. Typical form of an extrusion limit diagram.

The generation of heat during deformation is approximated using the assumption of adiabatic heating, which gives an upper bound estimate of the temperature rise. The temperature rise is considered in three areas of the extrusion: in main the deformation zone, ΔT1, at the interface between the container and the billet, ΔT2, and at the interface between the die land and the extrudate, ΔT3. These temperatures are estimated using the following expressions (developed by Stuwe [1] as quoted in [2]): ΔT1 = ΔT2 = ΔT3 =

σ ln R ρC p

(2a)

σ

vr lb ρ C p

4 3ρC p

k

σ

ve ld ρ C p

4 3 ρC p

k

(2b)

(2c)

where ρ is density (1.7 g/cm3), Cp is the heat capacity (1150 J/kg/K), vr is the ram speed, k is the thermal conductivity (130 W/m/K), ve is the extrusion speed and ld is the die land length. The stress was calculated using the hyperbolic sine law, an example of which is given below (modified from an equation derived in [1] for wrought AZ31 tested to a strain of 0.6): ⎛ ε& exp (147000 / RT ) ⎞ σ = 90 sinh −1 ⎜ 13 ⎟ 10 ⎝ ⎠

0.15

(3)

where T is in Kelvin, R is the gas constant and the strain rate is given by ε& = 9.6 R 0.6 vb / db . Ascertaining the temperature to use in determining the stress in this simple framework is, however, not so straight forward due to the coupling between stress and temperature. An appropriate value of stress was determined by averaging the stress calculated using the initial billet temperature and the stress calculated from the maximum temperature rise predicted assuming the deformation proceeds under the initial stress estimate. This average value of stress is then used to estimate ΔT1 for use in determining the maximum temperature achieved in the extrusion.

274 To generate an extrusion limit diagram it is necessary to perform a series of model simulations in which the speed is incrementally increased for a range of temperatures. For each simulation, the predicted loads and temperatures are compared with the maximum allowable values. Once these are exceeded, the conditions used in the simulation are recorded and then plotted to form an extrusion limit diagram. In the present work, a series of compression tests are employed to establish the influence of Al on hot strength. Laboratory extrusion tests are then used to ascertain the cracking limit and to tune the load predictions. The aim of the work is twofold: to verify the usefulness of the approach and to establish the influence of Al level on the extrusion limits of AZ alloys with Al contents below 3%.

Experimental Method Four AZ series alloys were cast with zinc and manganese levels held constant, and nominal aluminium levels of 0.7 wt%, 1.0 wt%, 1.5 wt%, and 2 wt%. Commercially obtained AZ31 was also tested as part of this study. Uniaxial compression tests were conducted on specimens which were 10mm in diameter and 15 mm long. The specimens were deformed at true strain rates of 0.1 s-1 and 1 s-1, at temperatures of 300 °C and 400 °C. Direct extrusion tests were carried out on a small scale laboratory rig (Figure 2a) designed for use in a 350 kN hydraulic press. The billet dimensions are 30 mm diameter by 20 mm in length. The billet, die and container are heated using electric heating elements. The temperature is measured using a thermocouple inserted into the container through a passage that terminates within ~2 mm of the container wall. For all of the present tests the extrudate was a round bar 5 mm in diameter. The tests were carried out at a constant speed, which was incrementally increased in subsequent tests until surface cracking was observed. An image of a cracked surface is contrasted with a normal finish in Figure 2b. Ram

Container Support Billet Die

Heating element

Insulation

a)

b)

Figure 2. a) Laboratory extrusion apparatus. b) Examples of extrudate surfaces obtained in the present work.

275

Results A typical series of stress-strain curves obtained using hot compression is presented in Figure 3a. The stress increases with increasing strain up to a peak after which work softening takes place at a rate that decreases with strain. This overall trend is common in materials displaying dynamic recrystallization [4, 5]. The level of the entire stress strain curve rises by ~25% when the Al level is increased from 0.7% to 3%. The uniform effect of Al on the stress-strain curve suggests that aluminium impacts on the deformation stress but not the initiation and progress of dynamic recrystallization. The flow stress at a strain of ~1.0, σ1, is plotted against the temperature corrected strain rate, Z = ε& exp (147000 / RT ) , in Figure 3b. Also included in this plot are predictions made using Equation 3 (short dashes), which was developed using compression and torsion tests carried out on fine grained wrought AZ31 [4]. It is clear that the present stresses, established using cast material, fall approximately 10 MPa higher. This difference may be ascribed to the fact that the steady state stress is approached more rapidly in fine grained samples due to the enhanced rate of dynamic recrystallization [6]. The following expression can be derived from the present data: ⎛ ε& exp (147000 / RT ) ⎞ σ = ( 85 + 6[ Al ]) sinh −1 ⎜ 13 ⎟ 10 ⎠ ⎝

0.15

(4)

Predictions made using this expression are shown in Figure 3b. It can be seen that the influence of Al is reproduced by the model. Due to the coarse grain size and the limited strain employed in the tests, the present data overestimate the actual steady state stress. For the present analysis, the pre-hyperbolic sine term in Equation 4 is adjusted to reconcile it with the same term in Equation 3, which is more representative of the steady state stress: ⎛ ε& exp (147000 / RT ) ⎞ σ = ( 72 + 6[ Al ]) sinh −1 ⎜ 13 ⎟ 10 ⎝ ⎠

(5)

200

100 80

150

1.5%Al 2%Al 3%Al

60 0.7%Al

σ1 (MPa)

Stress (MPa)

0.15

1%Al

100

0.7%Al 1%Al 1.5%Al 2%Al 3%Al 0.7% Al - predicted 3% Al - predicted 3%Al (fine grain)

40 50

20 400°C, 1.0 s

-1

0 1000000000

0

a)

0

0.2

0.4

Strain

0.6

0.8

1

1E+11

1E+13

1E+15

Z (s-1)

b)

Figure 3. Influence of Al on compressive flow stress. a) Stress-strain curves, b) Stress at a strain of 1 as a function of the temperature corrected strain rate, Z = ε& exp 147000 / RT .

(

)

276 The extrusion load was recorded in a series of experiments carried out under a range of different temperatures and speeds. The results are compared with predictions made using the approach presented above in Figure 4a. It can be seen that despite the many simplifications made, there is reasonable agreement between the model and experiment. The model tends to over-predict the stress at high temperatures (low stresses) and this may be due to the fact that, in most cases, these tests coincided with the occurrence of cracking. The lowering of the flow stress that occurs with decreasing Al level allows higher extrusion speeds for a given press capacity. That is, lower Al levels increase the low temperature extrusion limit. The extent of this increase in allowable extrusion speed was ascertained using the present model and the results of the simulations are shown in Figure 5b. Despite the moderate influence of Al on the flow stress, a speed increase of four fold is predicted to occur with lowering the Al level from 3% to 0.7%. 100

0.7%Al

500 0.7%Al 1%Al 1.5%Al 2%Al 3%Al

300

Ram speed (mms-1)

Measured load (kN)

400

1%Al

200

10 1.5%Al 2%Al 3%Al

100

1 250

0 0

100

200

300

400

500

300

350

400

450

Billet temperature (°C)

Predicted load (kN)

a)

b)

Figure 4. a) Actual extrusion loads as a function of predicted values. b) Influence of Al level on the low temperature (press capacity) extrusion limit.

To determine the influence of Al on the high temperature cracking limit, tests were carried out on each grade at increasing speed until cracks were observed. The conditions associated with cracking are shown in Figure 5a. The lines drawn through the data were fitted using the present model by iterating through different cracking limit temperatures. It is clear that the influence of Al on speed defined by this limit is greater than for the load limit. To illustrate the change in extrusion limit diagram achieved by reducing the Al content from 3% to 1%, both press and cracking limits are plotted in Figure 5b. 100 3%Al

1%Al

10

0.7,1%Al 1.5%Al 2%Al

1 300

400

500

Ram speed (mms-1)

Ram speed (mms-1)

100

600

1 250

Billet temperature (°C)

a)

10

3%Al

350

450

550

b)

Figure 5. a) Influence of Al level on the cracking limit. Filled symbols indicate successful extrusion, open symbols indicate cracking. b) Extrusion limit curves for Al levels of 1% and 3%.

277

Discussion The comparison between predicted and measured extrusion loads in Figure 4a shows reasonable agreement. We can therefore be confident in employing the model to predict the extrusion speed limit at low temperatures (load limit). In the present work, the position of the high temperature speed limit (cracking limit) was obtained experimentally. The model provides the shape of the limit curve. However, it should be noted that the assumption of adiabatic heating used in the model is a fairly crude approximation and will lead to increasing errors as the speed is decreased. For the present extrusion configuration, reducing the Al level from 3% to around 1% increased the maximum possible extrusion speed by 3-5 times. The precise increase achievable depends on how close to the limit it is possible to operate. Although the present work was carried out on a small laboratory press, the speed increases obtained are expected to be the same for commercial scale operations. In this regard it is interesting to note that the maximum speed predicted for the 3% Al alloy (AZ31) is approximately a third to a fifth of the maximum allowable speed, which is the speed at which a typical soft extrusion alloy (e.g. AA6061) would theoretically be able to be extruded. This difference tallies well with anecdotal comments in the literature with respect to the difference in extrusion speeds for the two metals [1].

Conclusion A simple analytical model was able to provide reasonable estimates of the extrusion load limit and fitting technique for the cracking limit. This approach was used to determine that decreasing the Al level from 3% to 1% increases the maximum extrusion speed by a factor of ~4 times.

Acknowledgments The help of John Vella and Graham Prior is gratefully acknowledged. The CRC for Cast Metals Manufacturing (CAST) is supported in part by the Australian Government Cooperative Research Centres Scheme.

References [1] F.E. Katrak, J.C. Agarawal, et al., Magnesium Technology 2000, ed. H.I. Kaplan et al., TMS, 351-354. [2] T. Sheppard: “Extrusion of Aluminium Alloys”, Kluwer, Dordrecht (1999), ch 2. [3] H.P. Stuwe: Metall., 22 (1968), 1197. [4] M.R. Barnett: Journal of Light Metals, 1 (2001), 167-177. [5] A.G. Beer and M.R. Barnett, Magnesium Technology 2002, TMS, Warrendale (2002), 193. [6] J.P. Sah, G.J. Richardson, and C.M. Sellars: Metal Science, 8 (1974, 325-331.

278

Comparison of Microstructure and Mechanical Properties of Indirect and Hydrostatic Extruded Magnesium Alloys J. Swiostek, J. Bohlen, D. Letzig, K.U. Kainer GKSS Research Center, Institute for Materials Research, D-21502 Geesthacht, Germany

Abstract The present study deals with the analysis of the microstructural development and the resulting mechanical properties during extrusion of magnesium alloys. Two extrusion processes – indirect and hydrostatic extrusion – will be compared corresponding to the influence of the aluminium content in AZ-alloys. The discussion will focus on the dependence of the profile microstructure and resulting mechanical properties on process parameters as well as on the extrusion process itself.

1

Introduction

Magnesium wrought alloys are of special interest today for the use as structural components due to the possibility to obtain even improved and more homogenous properties compared to cast components. Especially extrusion as form giving technology offers the possibility to produce such profiles that are long and/or have a thin-walled structure. While the process technology for the application on magnesium alloys is available today [1] the industrial use is not well established at present. More work is needed with respect to the understanding of the influence of extrusion as a thermo-mechanical treatment on magnesium alloys and therefore on the materials structural evolution. This is in order to set-up a well controlled and cost efficient process [2, 3]. Just little basic knowledge about important material aspects like microstructural evolution of the wrought material, its dependence on the processing and process parameters as well as the influence on mechanical properties is available. The present paper deals with the analysis of extruded alloys from the magnesium AZfamily containing a various amount of aluminum. While AZ31 is a well-known magnesium wrought alloy that is used as a correspondence material the influence of an increased content of aluminum in alloys AZ61 and AZ80 on the microstructural development and resulting mechanical properties will be discussed. It will be shown that the type of processing influences the evolution of the materials microstructure by using two extrusion processes. Beside of conventional indirect extrusion hydrostatic extrusion trials as a quite rarely applied technique were carried out which occurs to have a significant effect on resulting profiles.

279

2

Experimental

Commercial DC-cast billets were obtained in an as-cast condition from alloys AZ31 (Mg 2.9wt.% Al - 0.98wt.% Zn - 0.29wt.% Mn), AZ61 (Mg - 6.5wt.% Al - 0.99wt.% Zn 0.20wt.% Mn) and AZ80 (Mg - 8.5wt.% Al - 0.51wt.% Zn - 0.31wt.% Mn) in high-purity (HP) quality. All billets were machined to fit the extrusion presses and heat-treated for 12 h at 350 °C in case of AZ31 and AZ61 and at 385 °C in case of AZ80. The billets were air-cooled following ASM descriptions [4]. Extrusion trials were carried out using indirect extrusion and hydrostatic extrusion, respectively. The principles of the processes are shown in Fig. 1 for indirect extrusion and in Fig. 2 for hydrostatic extrusion.

Fig.1. Schematic indirect extrusion process.

Fig. 2. Schematic hydrostatic extrusion process.

During indirect extrusion the billet is put into a container and deformation is realized by pressing a hollow ram with a die in front of it into the billet. During hydrostatic extrusion the pressure is put onto the billet by using a hydrostatic medium. Thus pressure from a ram is put onto the medium and leads to hydrostatic pressure from all sides onto the billet to make it flow through the die. The difference between these two processes is the type of applied pressure on one hand and the die geometry on the other hand following the necessities of each process. Table 1 shows important parameters for the extrusion trials. For both processes and all three alloys the initial billet temperature was chosen to be 300 °C in order to make results comparable with respect to the aluminium content of the billet alloy and the type of the process. A rod-profile was extruded with an extrusion ratio of 1:23 in the case of indirect extrusion and 1:28 in the case of hydrostatic extrusion assuming that this is in good correspondence to each other. Due to the fact that only for the indirect extrusion process a continuous variation of the extrusion rate could be realized the slowest rate of the hydrostatic extrusion press was taken for convenience which refers to a profile speed of 8 m/min. However, since it occurred that indirect extrusion was limited in the possible profile speed for the alloys AZ61 and AZ80 the highest speed possible was applied in these cases. A microstructural analysis was carried out on polished and etched longitudinal sections using optical microscopy. For sample preparation an etching solution based on picric acid is used [5]. A commercial testing machine Zwick Z050 with a maximum applied force of 50 kN was used for tensile and compression tests at room temperature with samples fol-3 -1 lowing DIN 50125 and 50106. Strain rate was 10 s in all cases.

280 Tab. 1. Parameters of the extrusion trials.

Alloy

Extrusion method

Extrusion temperature [°C]

Extrusion ratio

Extrusion rate [m/min]

Die temperature [°C]

AZ31 AZ31 AZ61 AZ61 AZ80 AZ80

indirect hydrostatic indirect hydrostatic indirect hydrostatic

300 300 300 300 300 300

23 28 23 28 23 28

8 8 6 8 4 8

300 350 300 350 300 350

3

Results and Discussion

In all cases extrusion trials were carried out easily using process parameters like described in Table 1. A profile speed of 8 m/min was used since this has been the lowest speed accessible for hydrostatic extrusion. While during hydrostatic extrusion profiles with a smooth surface were obtained during indirect extrusion hot-shortness occurred for AZ61 as well as for AZ80. As an example micrographs of longitudinal sections of AZ80 in both – the hydrostatic (Fig. 3) and the indirect (Fig. 4) extruded – condition are shown for a profile speed of 8 m/min. It can be seen that cracks reach a depth of several millimetres for the indirect extruded rod while the hydrostatic one shows a smooth surface. A similar behaviour was found for AZ61.

Fig. 3: Hydrostatic extruded AZ80 (extrusion dir. →)

Fig. 4: Indirect extruded AZ80 (extrusion dir. →)

Therefore the extrusion speed has been reduced in these two specific cases leading to the maximum extrusion speed that allows profiles with a smooth surface. This speed was found to be 6 m/min for AZ61 and 4 m/min for AZ80. This is as a general effect in good correspondence to industrial parameters like shown in [6] where slightly lower speeds of 4 m/min for AZ61 and 2 m/min for AZ80 have been revealed.

281 Since hot-shortness is expected as the result of a heat-up due to friction between the billet and the container or the die, respectively, it can be stated that this heat-up is lower during hydrostatic extrusion than during indirect extrusion [7]. This allows to extrude alloys hydrostatically with a higher speed before hot cracking occurs. Longitudinal sections of the rods were used for preparing micrographs. They are shown in Figs. 5 - 10. The vertical direction of all micrographs represents the extrusion direction. In both rods from AZ31, indirectly and hydrostatically extruded, an inhomogeneous microstructure is found including grains with a wide range of sizes. Furthermore large grains with a long extension in extrusion direction occur. This effect is less pronounced for the hydrostatically extruded rod. These grains can be deleted by a further heat treatment. Thus it can be stated that the resulting microstructure in the as-extruded condition is only partly recrystallised. By increasing the aluminium content in alloys AZ61 and AZ80 a general improvement of the microstructure with respect to the homogeneity is observed. This is understood as an improvement of the recrystallisation behaviour during the dynamic thermo-mechanical treatment of extrusion.

Fig. 5. AZ31 hydrostatic extruded (extrusion dir. ↑).

Fig. 6. AZ31 indirect extruded (extrusion dir. ↑).

Fig. 7. AZ61 hydrostatic (extrusion dir. ↑).

Fig. 8. AZ61 indirect extruded (extrusion dir. ↑).

282

Fig. 9. AZ80 hydrostatic extruded (extrusion dir. ↑).

Fig. 10. AZ80 indirect extruded (extrusion dir. ↑).

The average grain size was determined using several micrographs from each alloy. The results are listed in Fig. 11. However, for both rods from AZ31, indirectly and hydrostatically extruded, only recrystallised regions have been observed and are compared to the results of AZ61 and AZ80. It is found that the resulting grain size is always lower after hydrostatic extrusion than after indirect extrusion. While the difference occurs to be extremely pronounced for the recrystallised regions in AZ31 the difference is decreasing with increasing content of aluminium. For hydrostatically extruded rods an increase of the grain size is observed from 4 μm for AZ31 to 9 μm for AZ80. Beside of the large grain size of 22 μm for indirect extruded AZ31 for AZ61 and AZ80 a comparable average grain size of 12–13 μm is found. The generally lower grain size of hydrostatically extruded rods is in good agreement to the occurrence of hot-shortness which has been related to a different heat-up during the extrusion processes. Assuming that the profile temperature is lower for hydrostatic extrusion than for indirect extrusion this would result in a lower grain size due to a different recrystallisation behaviour [8]. On the other hand the solidus temperature lowers with increasing aluminium content and therefore influences the recrystallisation behaviour. Figs. 12–15 show mechanical properties that have been obtained from tensile and compression tests at room temperature like described above. In Fig. 12 tensile (TYS) and compressive yield strength (CYS) are collected for both processes, in Fig. 13 the ultimate tensile strength (UTS) is shown. Fig. 14 shows the elongation to fracture. Corresponding to the finer grain structure of the hydrostatically extruded rods tensile and compressive yield strength occur to be higher than for the indirectly extruded rods. Elongation to fracture is nearly the same in all cases between 21–23% and no significant differences could be revealed with respect to the statistical accuracy of the measurement. It is of special interest to analyse the difference between tensile and compressive yield strength. In all cases that are shown in the present work it is found that the tensile yield strength is higher than the compressive one. This significant difference is generally known for extruded magnesium alloys [9] and related to a different activation of deformation mechanisms in tensile than in compression direction. Moreover this is based in the hexagonal lattice structure of magnesium alloys. However, this anisotropy is dependent on the alloy and the process used for extrusion. In Fig. 15 the difference TYS - CYS between tensile and compressive yield strength is shown as a representative value for the anisotropy and plotted as a function of the alloy for both processes. It can be seen that for hy-

283 drostatic extruded rods the anisotropy is lower than for indirectly extruded alloys which refers to the lower grain size after hydrostatic extrusion. The difference is strongly pronounced for AZ31 but decreases with increasing aluminium content by almost disappearing for AZ80.

Average grain size [µm]

30

300

indirect hydrostatic

25

250

YS [MPa]

20 15 10 5 0

150 100

2

3

4

5

6

7

Al content [wt.%]

8

9

0

10

at RT 2

4

5

6

7

Al content [wt.%]

30

Elongation [%]

350 300 250 200

20

4

5

6

7

Al content [wt.%]

8

9

15

10

Fig. 13. Ultimate tensile strength for extruded AZ alloys.

2

indirect hydrostatic

100 80 60 40 20

at RT 2

3

4

5

6

7

Al content [wt.%]

8

9

3

4

5

6

7

Al content [wt.%]

8

9

Fig. 14. Elongation for extruded AZ alloys.

120

0

10

at RT

at RT 3

9

25

150 2

8

indirect hydrostatic

indirect hydrostatic

400

100

3

Fig. 12. Yield strength for extruded AZ-alloys.

450

TYS - CYS [MPa]

TYS hydrostatic CYS hydrostatic TYS indirect CYS indirect

50

Fig. 11. Grain size for extruded AZ alloys.

UTS [MPa]

200

10

Fig. 15. Difference between TYS and CYS for extruded AZ alloys.

10

284

4

Summary

The different behavior of AZ-alloys during indirect and hydrostatic extrusion was analysed. It has been shown that hydrostatic extrusion offers the possibility to increase the extrusion speed especially for alloys with higher aluminum content by still avoiding hotshortness which easily occurs during indirect extrusion. It was discussed that profile temperatures stay lower for hydrostatic extrusion which not only influences the processing itself but also the resulting microstructure and mechanical properties. Thus a generally finer grain structure as well as increased yield behavior and strength were found for hydrostatically extruded rods. It was discussed that the anisotropy between tensile and compressive yield behavior is influenced by the average grain size and therefore occurs to be lower for hydrostatically extruded rods.

5

References

[1] K.B. Müller, Magnesium Technology 2002, edited by H.I. Kaplan, TMS 2002 [2] W.H. Sillekens, J. Bohlen, Proceedings of the 6th International Conference and Exhibition on Magnesium Alloys and Their Applications, 2003 (symposium M: Research Programs). [3] F. Lass, A. Mertz, E. Hombergsmeier, H. Lanzerath, P. Baumgart, P. Cordini, Proceedings of the 6th International Conference and Exhibition on Magnesium Alloys and Their Applications, 2003 (symposium M: Research Programs). [4] ASM Specialty Handbook: Magnesium and Magnesium Alloys (Eds. M.M. Avedesian, H. Baker), ASM International, 1999 [5] V. Kree, J. Bohlen, D. Letzig, K.U. Kainer, Praktische Metallographie, 2003, in press [6] G. Fischer, Documentation of the Conference „Magnesium im Fahrzeugbau“, IIR Deutschland GmbH, Sindelfingen, 2003 [7] W.H. Sillekens, J.A.F.M. Schade van Westrum, A.J. den Bakker, P.-J. Vet, Materials Science Forum 2003, Vols. 426-432 (Proceedings of THERMEC’ 2003), 629-636 [8] A. Beck, Magnesium und seine Legierungen, Julius Springer Verlag, Berlin, 1939 [9] J. Becker, G. Fischer in Magnesium – Eigenschaften, Anwendungen, Potenziale (Ed. K.U. Kainer), Wiley-VCH, Weinheim, 2002, 95-110

Improvements to the Rolling-Process in Order to Increase the Formability of Magnesium Wrought Alloys Friedrich-Wilhelm Bach*, Michael Rodman*, Mirko Schaper*, Arne Rossberg*, Eckart Doege**, Gerrit Kurz** * Institute of Materials Science (IW), University of Hanover, Germany ** Institute for Metal Forming and Metal Forming Machine Tools (IFUM), University of Hanover, Germany

1

Introduction

Due to their high specific mechanical properties, magnesium sheets offer a significant weight saving potential in modern vehicle constructions. To increase the range of applications e.g. in the automotive industry, deep-drawn magnesium sheets have to comply strict requirements regarding corrosion resistance, geometric tolerances, forming capacity and surface quality [1]. Furthermore the developers of a process-chain for the fabrication of magnesium sheets have to ensure cost effective and environmental friendly productionsteps in order to compete with alternative materials such as aluminum, modern steelalloys and plastics. The standard process-chain consists of slab-casting, rolling (incl. heat treatment), deepdrawing, joining, corrosion protection and enameling [2]. The present article covers new basic approaches to improve the rolling-process in order to suit the criteria mentioned above. The first part summarizes the research work conducted at the Institute of Materials Science (IW) at the University of Hanover concerning the rolling process. Besides basic parameter studies on thin magnesium strips, sample sheets were produced. The second part of the article covers the mechanical properties-testing and deep-drawing experiments of these sheets which took place at the Institute for Metal Forming and Metal Forming Machine Tools (IFUM) at the University of Hanover.

2

The rolling process

2.1

Sheet quality and rolling-parameters

Magnesium sheets for deep-drawing applications have to meet strict quality requirements which can be achieved through suitable rolling parameters. Table 1 gives a general overview of the influences of these parameters on the quality of magnesium sheets. The corrosion resistance depends strongly on the purity of the surfaces. In the heat treatment a growth of a MgOVMgO layer can be observed. Since this layer is unstable (Pilling-Bedworth-relation: < 1 ) in VMg contrary to aluminum, this has to be minimized. When lubricants are applied, it is important to examine if potentially harmful residues are rolled-in in the surface.

286 In order to achieve small geometric tolerances the rolling conditions have to be kept as homogenous as possible. This means among others that the temperature-drop between the front and rear of the rolling stock has to be minimized. For magnesium with its low effective heat capability the contact to the cold rolls and roller table can be unfavorable, especially for thin sheets. Therefore borders usually have to be cut and are returned to the melting furnaces. Further factors that influence the geometrical tolerances are the diameter- and concentricity tolerance of the rolls and the equable distribution of lubricants. Table 1. Influences of the rolling parameters on the sheet quality [3, 4, 6]

R pa oll r am ing ete req Qua rs u ir lit y em en ts

Heat treatment

Stable temperature conditions

Adequate strain

Rotation of the sheets

Suitable lubrication

Corrosion resistance (high purity)

++ 4

O

O

O

+++

Geometric tolerances

O

+++

++

++

+++ 3

+

++

+++

+

+

Formability for deep-drawing

+++ 1

+

++ 1

+++

+

Isotropic properties (decreased texture)

+

++

+

+++ 2

+

Roughness (low surface profile)

++ 4

O

+++

++

+++ 3

Border cracks

O - no influence +++ - high influence

1 - grain-refinement 2 - quasi-isotropic grain shapes

3 - lubricants decrease pressure 4 - formation of MgO-layer

Border cracks result from exceeding the tolerable tensions in the border-zones. This can result from inadequate strain (ε) and/or too low temperatures. The formation of textures is influenced by all rolling parameters and the heat-treatment of the specimens. For deep-drawing applications strong textures and the resulting anisotropy of the mechanical properties are counterproductive, hence a suitable, holistic rollingstrategy has to be applied. The surface-quality is evaluated by its purity and roughness. The accumulation of Mg/MgO-layers on the rolls resulting from fretting and the collection of residues of atmospheric particles or lubrication have to be avoided. Concerning the roughness, presently there is no standard for magnesium. In the enameling process for aluminum and steel sheets average peak-to-valley height (Ra) – values of 0.5-1.0 μm are acceptable. 2.2

Testing method

The rolling experiments were conducted on a twin-roll stand (manufacturer: Bühler, Type: Duo 125) with a diameter of the rolls of 129,8 mm and a rolling-speed of 0,13 m/s. The surface of the rolls was grinded and cleaned manually before each experiment. The result

287 ing surface of the rolls showed an average peak-to-valley height (Ra) in rolling direction (RD) of 0,65 μm and in transverse direction (TD) of 0,68 μm. In the presented experiments the following wrought-alloys were examined: MgAl3Zn1 (AZ31), MgAl2SE1 (AE21), MgMn1,2 (M1), MgZnSEZr (ZEK100) The dimensions of the specimen were: thin strips: length: 170 mm, width: 16-19, thickness 1,3 mm. sample sheets: length: 240 mm, width: 220 mm, thickness 1,0 mm. Due to the hexagonal lattice structure of magnesium the cold forming capacity is very limited. Above 200°C the ductility is increased by the activation of pyramidal slip-planes. Therefore hot-rolling cannot be avoided when forming magnesium wrought alloys. For the laboratory scale rolling experiments the feedstock was heated up to 400°C and rolled on heated cylinders with a temperature of T = 67 °C. It has to be mentioned that it is crucial to start the rolling process with suitable feedstock. The criteria that have to be met are a small, homogeneous grain size, high-purity condition and avoidance of micro and macro defects as well in the inner of the material as in the surfaces. For small and medium scale industrial production such feedstock is produced by slab-casting and milling of the outer layers, but it is also conceivable to adapt stripcasting techniques to magnesium. In the present research work (concerning the influence of the rolling-parameters) pre-rolled MgAl3Zn1 (AZ31) feedstock of s0 = 1.3 mm was used. For the production of sample sheets extruded profiles (60 x 170 x 6 mm) of the alloys mentioned above were rolled. Hence in both cases it was not necessary to break up the cast-microstructure before rolling. The basic influences of varied parameters in the heat-treatment and the temperature control have been published earlier [3]. Advanced studies are currently conducted or planned. 2.3

Application of lubricants and variation of the strain

The main aspect for the examination of these parameters is the avoidance of fretting on the rolls and thus decreasing the roughness on the sheets. When rolling magnesium without lubricants even at a low strain of 5%, after only a few reduction stages a layer of Mg/MgO on the rolls is clearly visible. This effect (fretting) rapidly builds up a layer of residues which cracks up (flittering) and is being rolled-in in the surface of the sheets, causing an increased roughness or even macro-defects. Besides constant cleaning (brushing) of the rolls it is possible to avoid fretting by application of suitable lubricants and rolling with moderate strains. The following lubricants were applied manually on the rolls: Table 2. List of examined lubricants (extract)

Product (1) Hycut® SW220 (2) MULTIDRAW® DRYLUBE C1 (3) MULTIDRAW® DRYLUBE G1 (4) Boron nitride

Manufacturer Main-component Oemeta Chemische Werke GmbH diester oil Zeller + Gmelin GmbH & Co hydrocarbonic compound of high molecular weight with polar additives Zeller + Gmelin GmbH & Co fully synthetic soap with polar additives Büro für angewandte Mineralogie, hexagonal boron nitride Dr. Stephan Rudolph powder Particle-size: 5-8 μm

288 The effectiveness of the lubricants is measured by the decrease of the coefficient of friction μ and the lowering of the rolling pressures Pw (measured rolling force dived by contact area). The lubricants were used as “lubricant in water”-solutions or suspensions with the most favorable concentration (as shown in previous publications [4]): (1) 5v%

(2) 5v%

(3) 20v%

(4a) 100%

(4b) 3m% in (1)

Lower concentrations show significantly less effectiveness, higher concentrations cause no increase and can cause instable rolling-conditions. Figure 1 shows the effectiveness of the selected lubricants in comparison to dry-rolling and the application of water. 450

0,45

Coefficient of friction

400

0,4

350

0,35

300

0,3

250

0,25

200

0,2

150

0,15

100

0,1

50 0

dry rolling

water

1

2

3

4a

Hycut® SW220

Multidraw® Drylube C1

Multidraw® Drylube G1

Boron nitride (BN 100%)

4b

Coefficient of friction µ [-]

Rolling pressure PW [MPa]

Rolling pressure

0,05 0

Hycut® SW220 + 3m% BN

Figure 1. Dependence of coefficient of friction and rolling pressure on lubrication

The significant decrease of both measured values for lubricant 1, 2 and 3 proves their general suitability. The effect of fretting and the growth of an Mg/MgO-layer on the rolls is reduced strongly and at certain strains it can be eliminated. Lubricant 4a (pure BN) decreases the rolling pressure excessively and can cause instable rolling conditions. The suspension 4b improves the results of lubricant 1 and thereby shows the possibility of using BN-powder as an additive to oil-in-water suspensions. The roughness of the sheets is a decisive factor considering the enameling of deepdrawn components. In the rolling process the initial value can set by variation of the rolling parameters. Figure 2 shows the influence of different lubrications on the average peak-to-valley height Ra – values in RD and TD.

289 1,6

RD

Average peak-to-valley height Ra [µm]

1,4

TD

1,2 1,0 0,8 0,6 0,4 0,2

1

0,0

grinded feedstock

dry rolling

water

2

Hycut® SW220

3

Multidraw® Multidraw® Drylube C1 Drylube G1

4a

4b

Boron nitride (BN 100%)

Hycut® SW220 + 3m% BN

Figure 2. Ra in RD and TD in variation of the lubricant

The roughness was measured using a contactless laser-optic tester (Rodenstock RM 600) in accordance with DIN 4762. The results show higher Ra-values in TD than in RD and that formed specimens generally have a lower roughness than the feedstock. The only exception is registered when instable rolling conditions occur (lubricant 4a). Further it can be noticed that all lubricants increase the roughness in comparison to dry-rolling. This leads to the conclusion that the smoothing of the surface depends strongly on the rollingpressure, which is lowered by the lubricants. To prove this thesis extensive experiments at different strains (5 % < ε < 26 %) have been conducted. Figure 3 shows the effectiveness of lubricant 1 at different strains as the reduction of the rolling-pressure in comparison to dry-rolling. 350

Rolling pressure PW [MPa]

dry rolling

Hycut® SW220

300

-7,1%

250

-4,5% 200

-17,7%

-6,9% 150

-5,7%

100 50 0

ε = 5%

ε = 10%

ε = 15%

Figure 3. Effectiveness of lubrication at different strains

ε = 20%

ε = 26%

290 The highest decrease in Pw is observed for a strain of 15%. This effect is interpreted with two opposed effects: a) up to a strain of 15% the effectiveness of the hydraulic wedge increases due to the separation of the friction partners and b) above 15% the lubricationfilm begins to break up and the rolling pressure grows over proportionally. This barrier for the strain is also visible in form of growing residue-layers on the rolls at a strain above 15%, which is 10% higher than for dry rolling. 1,0

Average peak-to-valley height Ra [µm]

0,9

RD

TD

0,8 0,7 0,6 0,5 0,4 0,3 0,2 0,1 0,0

dry

Hycut® SW220

ε = 5%

dry

Hycut® SW220

ε = 10%

dry

Hycut® SW220

ε = 15%

dry

Hycut® SW220

ε = 20%

dry

Hycut® SW220

ε = 26%

Figure 4. Roughness for dry and lubricated specimens at different strains

In figure 4 the roughness in RD and TD is shown in dependence of the strain with and without the presence of lubricant 1. Up to a strain of ε = 20% the roughness of the specimen decreases for dry and lubricated rolling in RD and TD. At higher strain it stays constant. In rolling direction RD the value is as generally known always lower than in transverse direction TD, but at a strain of ε = 15% Ra in RD and TD show the smallest difference which points to homogenous rolling-conditions and correlates with the highest effectiveness of the lubricant at this point. In conclusion it can be said that for the analyzed rolling-process the best compromise between low-rolling-pressure (to avoid fretting) and low roughness (for the enameling) can be achieved with suitable lubricants at strains around ε = 15%. At higher strains fretting, which can cause macro-defects on the sheets, can not be avoided completely only by lubrication. An additional brushing system seems to be appropriate. 2.4

Rotation of the sheets

The main arguments for rotating the specimens between the rolling stages are the minimization of the thickness-tolerances, the homogenization of the grain-shape and size and the decreased texturing-intensity. When rolling uni-directional the fluctuation in thickness in RD is twice as high as in TD due to lower temperatures in the rear of the rolling stock.

291 Therefore (if real-time adjustment is not possible) the rotation by 90° and 180° after certain stages is an industrially-meaningful modification. The thickness fluctuation in RD and TD of these fine-rolled and rotated sheets was between 0,97 – 1,08 mm which is nearly half of uni-directionally rolled sheets. a) standard rolling (0 0 2) - plane

b) rotated rolling (0 0 2) - plane

AE21, 5th pass, ε = 10,5 %, d = 5,2 mm AE21, 15th pass, ε = 9,6 %, d = 1,7 mm Figure 5. Textures in the (0 0 2) - plane for standard- and rotated-rolled MgAl2SE1 (AE21)samples.

Figure 5 shows the textures in the (0 0 2 plane) of a) uni-directionally-rolled and b) rotationally-rolled sheets. The left half displays a strong orientation in the crystal-structure resulting from the rolling stages. On the right side this former structure is still visible but it is overlaid by a new perpendicular orientation resulting from turning the sheets by 90°. Furthermore the intensity of the textures is decreased by 1/3.

3

The deep-drawing process

3.1

Mechanical properties

In order to evaluate the suitability of the sample sheets for deep drawing processes, flow stress/logarithmic strain-diagrams (kf vs φ) have been recorded in tensile tests (EN 10002/5) at T = 20 ° C and T = 200 °C. All specimens were heat treated after rolling, in order to ensure adequate strain for the following forming steps. Figure 6 shows that generally the achievable strain grows and the yield-stress decreases when reaching the activation-temperature of the pyramidal-slip (depending on the alloy around T = 200 °C). This material behavior corresponds with the results in previous publications [8]. At room temperature MgAl2SE1 (AE21) and MgMn (M1) show poorer forming behavior in comparison with MgZnSEZr (ZEK100) and standard AlMg3Zn1 (AZ31). At a temperature of T = 200 °C the magnesium alloy MgMn (M1) shows the best formability.

292 350

MgAl3Zn1 (AZ31) (T = 20 °C)

Flow stress kf [MPa]

300

ε = 0.004 s-1

MgAl2SE1 (AE21) (T = 20 °C)

250

MgMn (M1) (T = 20 °C) MgZn1SEZr (ZEK100) (T = 20 °C)

200

MgAl3Zn1 (AZ31) (T = 200 °C)

150

MgMn (M1) (T = 200 °C)

100

MgAl2SE1 (AE21) (T = 200 °C)

50

MgZn1SEZr (ZEK100) (T = 200 °C)

0 0

0.1

0.2

0.3

0.4

0.5

Logarithmic strain ϕ [ - ] Figure 6. Flow stress/logarithmic strain-diagrams for heat treated MgAl3Zn1 (AZ31), MgAl2SE1 (AE21), MgZnSEZr (ZEK100), MgMn (M1) at T = 20 °C and T = 200 °C

3.2

Deep-drawing results

When interpreting results from uni-directional tensile-tests it has to be considered that the transferability on the deep drawing behavior with resulting multi-axis-tension is limited. Therefore standard cup-deep-drawing test have been conducted. The cup-deep-drawing tests show a significant influence on the composition of the alloys, as shown in previous publications, too [8], [9]. limit drawing ratio

Limit drawing ratio ß 0,max [-]

3,2 3,0

1

Ø = 50 mm rP = 8 mm rD = 6 mm

2,8 2,6

2

ß0, max =

1: punch diameter d0 2: maximum blank diameter D0,max

2,4 2,2 2,0

MgZn1SEZr (ZEK100)

1,8

MgAl2SE1 (AE21)

1,6 1,4 1,2 1,0 125

150

175

200

225

250

275

300

325

Temperature T [°C] Figure 7. Limit-drawing-ratios of AE21 and ZEK100 in variation of the temperature

350

D0,max d0

293 Figure 7 shows the limit-drawing-ratio (β0,max) of MgAl2SE1 (AE21) at temperatures between T = 150 °C and T = 325 °C. It is visible that at T = 300 °C the limit drawing ratio β0,max is optimized (β0,max = 2,2) for this alloy. In contradiction to the results of the tensile tests, the limit-drawing-ratio for MgZnSEZr (ZEK100) at T = 225 °C is significantly higher (β0,max = 3,0) than for MgAl2SE1 (AE21). This incongruity can be understood when taking into consideration the greater intensity of the rolling-textures for MgAl2SE1 (AE21).

4

Discussion and evaluation

The rolling-experiments show that improved surface quality can be achieved when using suitable lubricants at a strain around ε = 15%. Intolerable effects such as fretting and flittering resulting from residue-layers on the rolls can be minimized. Nevertheless brushing of the rolls seems sensible for mass production. The roughness of the sheets can be adjusted by an adequate strain in the last rolling stage. The geometric tolerances of the sheets are improved by rotating the sheets in the rolling process. A further effect is the partial compensation of uni-axial textures. The tensile tests show a significant influence of the composition of the alloys and the temperature on the forming behavior. The cupdrawing-test for MgZnSEZr (ZEK100) proves that a limit-drawing-ratio of β0,max = 3,0 at T = 225 °C, which is higher than for aluminum alloys, is possible. Coming experiments will determine the influences of the temperature-management, different rolling-textures and the variation of the roughness of the rolls in order to further improve geometric tolerances, forming behavior and surface quality. The deep-drawing experiments have to be conducted for complex shapes and plant-modifications (e.g. hydro-forming) are to be evaluated.

5

References

[1] S. Schumann, H. Friedrich, Materials Science Forum, Vols. 419-422 (2003), pp. 51-56. [2] P. Juchmann, Appropriate process chain for ultralight components out of Magnesium sheet metal for transportation (ULM), Materials Week 2002, Munich. [3] H. Haferkamp, J. Zilberg, M. Rodman, M. Niemeyer, J. Weber, U. Holzkamp, Theorie und Praxis der Metallurgie , N5, Heft19, 2000, ISSN 1028-2335, pp. 23-24. [4] Fr.-W. Bach, M. Rodman, J. Zilberg, A. Roßberg, Metall, 57. Jahrgang, ISSN 00260746, 5/2003, pp. 294-298. [5] Fr.-W. Bach, M. Rodman, M. Schaper, J. Weber, J. Zilberg, Aluminum, planned in 2003. [6] F. Kaiser, D. Letzig, J. Bohlen, A. Styczynski, Ch. Hartig, K.U. Kainer, Material Science Forum Vols. 419-422 (2203, pp. 315-320. [7] H. Haferkamp, E. Doege, M. Schaper, Ch. Jaschik, M. Rodman, G. Kurz, Metall, 56. Jahrgang, ISSN 0026-0746, 12/2002, pp. 712-716. [8] K. Dröder, Untersuchungen zum Umformen von Feinblechen aus Magnesiumknetlegierungen. Dissertation, Universität Hannover, 1999. [9] E. Doege, W. Sebastian, K. Dröder, G. Kurz, Increased Formability of Mg-Sheets using Temperature Controlled Deep Drawing Tools, in Innovations in Processing and Manufacturing of Sheet Materials, TMS Annual Meeting, 11.-15.02.2001, p. 53-60

The Effect of Strain-Hardening and Process Parameters on the Forgeability of Wrought Magnesium Alloys Johnson C.F. Chan* †, C.J. Tay†, M.S. Yong* * Singapore Institute of Manufacturing Technology † National University of Singapore Email: [email protected]

Abstract Magnesium alloys generally are known to be brittle and have poor deformability at room temperature. This can be attributed to its hexagonal close packed lattice structure. Commercial magnesium alloys of AZ31 and AZ61 are studied and tested. Uniaxial compression tests were conducted under various temperatures to establish the strain-stress, straintemperature relationships and its influences on the microstructure. The objective of this investigation is to establish the relationship between the deformation temperature vs. the strain-hardening, n value and strength coefficient, K value. The paper also attempts to establish a range of optimal temperature for the forging of wrought magnesium alloys.

1

Introduction

One of the main advantages of magnesium alloys is its inherent lightweight property. Among several industries applications, one of the distinctive trends is towards the weight reduction in vehicles construction [1]. Although the magnesium alloys have several advantages over many metals, the limitations in plastic deformation capabilities with lower working temperatures restrict a wide spread use to the forging and forming processes. This is attributed to the hexagonal crystal packed structure of the alloys, where it limits this structure to a single slip planes. However recent findings show the feasibilities in working with magnesium alloys at higher temperature and a number of researchers [2][3][4] have successfully forged and formed the magnesium alloys. The improved ductility with magnesium alloys at higher temperatures enable the activation of slip planes and thus, improves plastic deformation. Although magnesium alloys may possess greater forgeability at higher temperatures, it is still essential to analyse on the suitable working temperatures and forces being applied. In this present work, the strength coefficient (K) and work hardening (n) factors will be studied to enhance understandings on the forgeability for these alloys. Investigation on the optimum temperatures for the wrought magnesium alloys is also been conducted to study on the influence of temperature correlation between the K and n factors.

295

2

Experimental Procedure

AZ31 and AZ61 wrought magnesium billets were tested as accordance to ASTM uniaxial compression standards [5] and their composition contents are shown in Table 1. The process parameters considered include temperatures ranging from ambient (22 ºC) to 300 ºC at an increment of 50 ºC, and compressive loads ranging from 55 KN to 60 KN with a constant strain rate at 2.54 s-1. The true stress-strain results are subsequently obtained from the compression testing at various temperature settings and provide the fundamental studies on plastic deformation behavior for the wrought alloys. The K and n factors are established from the true stress-strain curves. Table 1. Composition Contents for AZ31 and AZ61

Wt% AZ31 AZ61 2.1

Al 2.8 6.3

Zn 0.9 0.73

Mn 0.24 0.19

Fe 0.004 0.003

Si <0.01 <0.01

Cu <0.01 <0.01

Ni >0.001 <0.001

Ca <0.01 <0.01

Methods to Obtaining K and n Factors

The log true stress – log true strain plot is required when determining the strength coefficient and strain-hardening factor. The Hollomon analysis is the most commonly used empirical stress-strain analysis method. This analysis assumes flow curve of many metals in the uniform plastic deformation region and is simply expressed by: σ = K εn

(1)

where n is the strain-hardening exponent and K is the strength coefficient. A log-log plot of the true stress and true strain up to maximum load will results in a straight-line if is satisfied by the data. The linear slope of this line is n and K is the true stress at ε = 1.0. The n exponent may have values from n = 0 (perfectly plastic solid to n=1 (elastic solid). Most metals have n ranging from 0.10 to 0.50. The strain hardening exponent also measures the maximum uniform strain of the specimen and usually represent the maximum practical elongation in engineering application, so an increase in the stain hardening increases the strain to failure.

3

Results and Discussion

The true stress-strain curves obtained for both AZ31 and AZ61 wrought magnesium alloys are shown in figure 1. Table 2 displays the uniaxial compression results including the required forging pressure, true strain values and elongation percentages for both the wrought magnesium alloys. Figure 2 shows the true strain values plotted at different testing temperatures for both the wrought alloys. The true strain at the lower temperatures compression was insignificant due to the effect of fracturing by the wrought magnesium alloys. Notably, these alloys exhibit poor plastic deformation between ambient to 100 ºC. This low ductility effect is mainly due to the inactive slip planes within the hexagonal crystal packed structure for these wrought alloys

296 tested at lower temperatures. Further analysis shows that higher flow stresses are needed to compress the material at ambient and 100 ºC conditions, where material fracture are expected to occur at true strain of around 0.17 and 0.27 for both the wrought alloys. See Table 2 for more details.

Figure 1. Uniaxial compression true stress vs. strain curves for AZ31 and AZ61 magnesium alloys at various temperatures. True Strain 0.7

Table 2. Experimental results tabulated on final pressure, true strains and elongation % on AZ31 and AZ61 for all temperatures. *(F) denotes material failure

AZ31

0.6

AZ61

Material Temperature Final forging True strain Elongation ºC pressure (MPa) (%)

0.5 0.4 0.3

AZ31

22 100 150 200 250 300

440.97 397.99 188.26 116.21 89.52 79.53

0.161 (F) 0.27 (F) 0.508 0.608 0.649 0.65

9.47 (F) 14 (F) 53.36 70.27 77.6 80.57

AZ61

22 100 150 200 250 300

441.04 427.36 234.73 119.82 84.95 77.86

0.168 (F) 0.25 (F) 0.463 0.583 0.65 0.65

8.8 (F) 11.6 (F) 40.85 71.07 79.07 80.79

0.2 0.1 0

100

200

300

400

Temperatures ºC

Figure 2. Experimental results showing the Strain limits obtained at different temperature increments.

Further tests conducted with higher temperatures displayed a large plastic deformation improvement for both the wrought magnesium alloys. The low strain and fracturing effect achieved at the earlier stages were represented by the large increased in deformation strain -1 values. Figure 2 shows that true strain at 0.65 s was achieved for both the magnesium alloys and is equivalent to the superior 80% plastic deformation for both the wrought magnesium alloys. The improved ductility for the wrought alloys can also be observed by the lower true compressive stresses. The stresses decreased significantly with increasing temperatures and strain values. The results on the final pressure are shown in Table 2 and

297 the data represents a substantial decrease in the compressive stress as the deformation temperature operates at 200 ºC and above. The improved ductility is attributed to the flow softening effect within the materials, largely due to the thermally activated motion of dislocations, with contribution from the additional of slip planes within the wrought alloys. The effect of flow softening is noted at the onset for experiments conducted at 150 ºC and above. Figure 2 shows the large differences in true stains between the 100 ºC and 150 ºC test results. Table 3. The grain sizes at various temperature regimes measured from the micrograph analysis. The measurement for the grain sizes are averaged from 3 specimens tested with similar testing conditions.

Materials

Origin

100 ºC

150 ºC

200 ºC

250 ºC

300 ºC

AZ31 AZ61

12-14 μm 13-15 μm

9-10 μm 8-9 μm

5-6 μm 5-6 μm

4-5 μm 4-5 μm

2-3 μm 2-3 μm

2-5 μm 3-5 μm

Table 3 displays the average grain sizes measured from the micrographs for both the wrought magnesium alloys deformed between 100 ºC to 300 ºC. Lower ductility of the wrought alloys is evident from the large grain sizes obtained for parts deformed at lower temperatures and both alloys fracture during the early compressive stages. At 150 ºC onwards, the grain sizes begin to decrease after every successive temperature increment, thus showing superior plastic deformation. At 300 ºC, observation show that there is a slight increase on the grain sizes and the formation of larger grains may be due to the onset of dynamic recrystallization (DRX) for both the magnesium alloys [6][7] for this working temperature. It is notable that DRX contributes to the material softening and the improved ductility for AZ31 and AZ61 alloys. 3.1

K and n values

Up to date, very little works has been conducted to study the phenomenon of strain hardening exponent, (n) for magnesium alloys [7]. Research findings are based on uniaxial tension testing [8][9] and common materials used are aluminum-magnesium based alloys. Figure 3 shows the various strain-hardening exponents (n) achieved at different temperature increments for both the alloys. Table 4 displays the K and n values fitted with the Hollomon equation. The stain-hardening exponent, n is obtained from the logarithmic true stress-strain plot. Results showing that the n values decrease with temperature increment [10]. An AZ61 alloy exhibits a higher strain hardening value at 0.89 during the 150 ºC testing. At 150 ºC, the AZ61 material is not as sensitive to isothermal forging effect. This effect is further evident from the tabulated results in table 2, where only 40.82% of elongation is achieved. On the other hand, AZ31 exhibits lower strain-hardening value of 0.68, therefore exhibiting better ductility and material softening. Result also shows that a higher percentage elongation is achieved at 53.36% at 150 ºC. Hence, the effect of flow softening was more pronounced with the AZ31 specimens at this temperature. AZ31 and AZ61 alloys exhibit a distinctive decrease in its strain-hardening value when the next temperature increment is reached. At 200 ºC, the lower strain-hardening value between 0.22 and 0.24 were achieved for both AZ31 and AZ61 respectively, thus demonstrating the largest material softening effect. In accordance to bibliographic results

298 [11][12], the studies on magnesium alloys have shown that up to 225 ºC, slip occurs readily on the basal planes (0001) of the hexagonal crystal structures. This explains the operative slip planes in favorably orientated in relation to the direction of stress and thus correlating superiorly to the deformation capability. Experiments with further temperature increment till 300 ºC show the lowering of strainhardening values slightly for both the wrought alloys. This is more obvious with the AZ61 specimens where it exhibits a further decrease in its n value at 0.15. The decrease in the n value suggested further flow stress softening. The reading from Table 2 provides evident that lesser flow stress (77.86 MPa) is needed to compress the AZ61 specimen. Table 4. The strength coefficient (K) and strainhardening factor (n) fitting with the flow curve equation for temperature ranging from 150 ºC to 300 ºC.

Temp AZ31 – n ºC σ=Kε 150 200 250 300

σ = 177.86 ε0.68 σ = 144.40 ε0.22 σ = 83.62 ε0.19 σ = 60.19 ε0.198

AZ61 – σ = K εn σ = 151.99 ε0.89 σ = 86.54 ε0.24 σ = 70.74 ε0.17 σ = 66.85 ε0.15

Figure 3. Results for the strain-hardening values plotted with temperatures ranging from 150 ºC to 300 ºC.

4

Conclusion

The following conclusions can be drawn: 1. At ambient (22 ºC) and 100 ºC, the magnesium alloys are found to possess poor deformability and both AZ31 and AZ61 fracture at lower strains. This is mainly attributed to its hexagonal crystal packed structure having a single slip plane at lower temperature. 2. The limit strain increases and lower stress is used to deform AZ31 and AZ61 alloys at 150 ºC and above. Results from the micrographs measurement also show smaller grain sizes being formed at 150 ºC and above when compared with its original grains. 3. The strain hardening exponent, n decreases with increasing temperatures. The decrease in n values, exhibit material softening and improved ductility. This is evident from the distinctive plastic deformation for AZ31 and AZ61 alloys. 4. At 200 ºC, the strain-hardening exponent exhibits large decrease due to material softening. It is therefore evident that true plastic localization for both the wrought alloys occurs near this temperature range. 5. The Hollomon analysis shown in Table 4 explains the stress-strain relationship and material behavior for AZ31 and AZ61 alloys. Result from the stress-strain curves shows that the strain-hardening effect is dependent on the temperatures and plastic strain range.

299

References [1] Vohringer, K.D., Advanced Technology of Plasticity, Springer, Berlin, 1999, pg 3. [2] Johnson, C.F. Chan., Yong, M.S., Tay, C.J., Shang, H.M., “The Influence of Process Parameters on Forged Magnesium Alloys”, ICAMP 2002 Grand Hyatt Singapore [3] Yong, M.S., Hu, B.H., Enggalhardjo, M. and Choy, C.M., “Development of Warm Forming Technique to Produce Thin Wall Magnesium Components”, 10th International Conf. on Precision Engineering (ICPE), 18-20 July 2001, Yokohama, Japan. [4] Shehata. F., Painter, M.J., Pearce. R., “Warm Forming of Aluminum and Magnesium Alloy Sheet”, Journal of Mechanical Working Technology, 2 (1978) 279-290. [5] “Annual Book of ASTM Standards”, 1998, Section 3, E 9-89a [6] Gehrmann, R., Gottstein, G., in: J. Szpunar (Ed.), Proceedings of the 12th Intl Conf. Texture of Materials. (ICOTOM 12), NRC Research Press, 1999, 665 [7] Yu, K., Li, W.X., Zhao, J., Ma, Z.Q., Wang, R.C., “Plastic deformation behaviors of a Mg-Ce-Zn-Zr alloy” Scripta Materialia 48 (2003) 1319-1323 [8] Shuaib, A.N., “ Mechanical properties of Al-2.5 Mg-0.1 Mn-Si-Cr-Fe alloys”, Materials and Design 23 (2002) 181-187 [9] Tetsuo Naka, Torikai, G., Hino, R., Yoshida, F., “The effects of temperature and forming speed on the forming limits diagram for type 5083 aluminum-magnesium alloy sheet” J.Mat Proc.Tech. 113 (2002) pg 648-653. [10] Martin, E., Forn, A., Nogue, R., “Strain hardening behavior and temperature effect on Al-2124/SiCp” J.Mat. Proc.Technology (2003) [11] Emley, E.F., “Principles of Magnesium Technology”, Pergamon Press, Oxford, New York, 1996, 231-241 [12] Chapman, J.A., PhD. Thesis, University of Birmingham, 1963

Superplasticity in Magnesium Alloys Processed by Different-Roll-Speeds Rolling Shangli Dong1, 2, Tsunemichi Imai1, Ichinori Shigemastu1, Naobumi Saito1 1

Institute for Structural and Engineering Materials, National Institute of Advanced Industrial Science & Technology (AIST), Nagoya, 463-8560, Japan; 2 Harbin Institute of Technology, Harbin, 150001, China

1

Introduction

As a light weight structure material candidate, magnesium and its alloys are increasedly applied in transportation, electronic and aerospace industries, and such trend will continue within the following 10 to 20 years [1, 2]. However, most currently used magnesium alloys are cast alloys due to the limited ductility and poor workability of magnesium at ambient or low temperatures [3, 4]. It is no doubt that high performance wrought magnesium alloys developed economically by bulk working techniques such as forging, rolling and extruding, would be essential to widespread application of magnesium alloys. Rolling is an important process for sheet production of metals and their alloys. The different-roll-speeds rolling (DRSR) process is an asymmetric rolling, in which the rotation speeds of upper roll and lower roll are different. Recent asymmetric rolling experiments carried out on aluminum alloys showed that intense shear deformation occurred uniformly through the sheet thickness and resulted in fine-grain with high angle boundaries [5-7]. Introduction of DRSR technique to magnesium sheet fabrication is consequently thought to produce grain refinement and mechanical property enhancement by the intense shear strain during rolling operation. Superplasticity, in metals, ceramics and intermetallics, is being utilized to form parts in many industry areas, which would provide producer with a suitable forming method [8,9]. Attaining and revealing the potential superplasticity in magnesium alloys are also important to the development and application magnesium alloys. Furthermore, understanding the effects of process on superplasticity of wrought magnesium alloys will give beneficial information for process optimization and bulk workability improvement of the alloys. Two primary attempts were made in present investigation. The first is to produce Mg-AlMn based (AM60) and Mg-Al-Zn based (AZ31 and AZ91) alloys sheet from cast billet using the DRSR process. Secondly, evaluation of superplasticity and superplastic blow forming were conducted in the as-rolled alloys. Relationship of superplastic characteristics to microstructure features and intense strain produced by DRSR has been discussed, and the favorable rolling procedure to obtain superplasticity in the examined alloys is suggested.

2

Experimental

The start magnesium alloy billets with size of 10×60×80mm for rolling were cut from AM60, AZ91 and AZ31 alloy cast ingots. DRSR was conducted with a mill, of which the

301 rotation speed ratio of the upper roll and lower roll can be varied from 1 to 1.364 by using the gear shift system while the diameters of both rolls are 160 mm. The rotation speed of the lower roll is kept as a constant of 10 m/min, and rolls can also be heated up to 300 °C. The investigated alloys were finally rolled to sheet materials with final thickness of 1.3 mm at specified temperatures for each given alloy, within the billet temperature range of 300 °C to 400 °C. The reduction ratio in thickness was selected as 5% per pass, and the ratio of 10% was also tried for somewhat comparison. Uniaxial tension test was applied to evaluate superplasticity at two initial strain rates of 1.4 ×10-3 and 1.4 × 10-2 s-1, and at temperatures from 300 °C to 450 °C. Tensile specimens with a gauge length of 6 mm and with the gage axis parallel to the rolling direction were cut from the as-rolled sheets without any heat-treatment. Tensile properties of the as-rolled sheet at room temperature was carried out with the crosshead speed of 2 mm/min. Constant pressure superplastic blow forming test was performed at elevated temperature on the as-rolled AM60 alloy using disk specimen of 60 mm in diameter. The Ar gas pressure for forming operation could be managed by a computer controlled automatic feedback system.

3

Results and Discussions

3.1

Investigation on Rolling Process

Several rolling practices carried out on AM60, AZ31 and AZ91 alloys are listed in Table 1. The rolls were not lubricated at room temperature while lubricated at some elevated temperature as indicated in the table. It was revealed that careful control of rolling conditions is important for magnesium sheet production. Any crack produced during rolling must be cut in order to get the final sheet free of cracks. The rolling direction and billet position were kept the same in all rolling passes to obtain more effective grain refinement suggested by asymmetric rolling research on aluminum alloys [5, 6]. AM60 and AZ91 sheets rolled under conditions of 400 °C/ R.T. /5% and 400 °C/100 °C/5 % are herewith selected for examination, while the sheets rolled at 300 °C/ R.T. /5% and 300 °C/100 °C/5 % for AZ31 alloy. The tensile properties of them at room temperature are shown in Table 2. The properties are favorable for each alloy and seem sufficient to meet the requirements of most industrial application. Table 1. Rolling process record of the investigated Mg alloys

Alloy AM60

AZ91

AZ31

Billet temperature 350 °C 400 °C 400 °C 400 °C 350 °C 400 °C 400 °C 400 °C 400 °C 300 °C 300 °C

Rolls temperature R. T. 100 °C R.T 100 °C R. T. 100 °C R. T. 100 °C 100 °C 100 °C R. T.

Rolling Reduction 5% 10% 5% 5% 5% 10% 5% 5% 5% 5% 5%

Remark Surface crack Edge and surface crack, large Edge crack, small No crack/Lubricated Surface/edge crack, very large Edge and surface crack, large Edge crack, small No crack/Lubricated No crack/Lubricated No crack/Lubricated Edge crack, small

302 Table 2. Tensile properties of the as-rolled Mg alloy sheets at room temperature

Alloy AM60 AZ91 AZ31

Ultimate strength 0.2% proof strength 290 MPa 193 MPa 294 MPa 196 MPa 319 MPa 200 MPa 307 MPa 197 MPa 248 MPa 185 MPa 254 MPa 191 MPa

Elongation 8.5 % 8.7 % 9.3 % 9.9 % 11.3 % 12.0 %

Rolling conditions 400 °C/R.T./5% * 400 °C/100 °C/5% 400 °C/R.T./5% 400 °C/100 °C/5% 300 °C/R.T./5% 300 °C/100 °C/5%

* 400 °C/R.T./5%denotes the billet temperature, rolls temperature and reduction ratio per pass, respectively.

3.2

Superplasticity Evaluation in AM60 Alloy

The total elongation and flow stress via testing temperatures of the AM60 sheets rolled at 400 °C with rolls temperature of R. T. and 100 °C are shown in Figures 1 and 2, respectively. It can be found that lower roll temperature is beneficial to increase elongation and reduce flow stress. The as-rolled AM60 alloy exhibits larger elongation at or over 400 °C at strain rate of 0.00141/s, which belongs to conventional superplasticity. No remarkable high strain rare superplasticity has been detected under the present rolling and tension testing conditions. 400

80

200

100

350

375

400

425

450

(b)

Strain rate= 0.014 1/s Strain rate= 0.0014 1/s

60

40

20

Strain rate= 0.014 1/s Strain rate= 0.0014 1/s 0 325

AM60 Alloy Rolls temperature= R.T. o Billet temperature= 400 C

(a)

Flow Stress, MPa

Total Elongation, %

300

100 AM60 Alloy Rolls temperature= R.T. o Billet temperature= 400 C

0 325

475

350

375

o

400

425

450

475

o

Temperature, C

Temperature, C

Figure 1. Total elongation (a) and flow stress (b) vs. testing temperature in the AM60 alloy rolled at 400 °C with the rolls kept at room temperature. 400 (a)

80 Flow Stress, MPa

Total Elongation, %

300

100 AM60 Alloy o Rolls temperature= 100 C o Billet temperature= 400 C

200

100

0 325

350

375

400

425 o

Temperature, C

450

475

(b)

60

40

20

Strain rate= 0.014 1/s Strain rate= 0.0014 1/s

AM60 Alloy o Rolls temperature= 100 C o Billet temperature= 400 C

0 325

Strain rate= 0.014 1/s Strain rate= 0.0014 1/s 350

375

400

425

450

475

o

Temperature, C

Figure 2. Total elongation (a) and flow stress (b) vs. testing temperature in the AM60 alloy rolled at 400 °C with the rolls heated at 100 °C.

303 3.3

Superplasticity Evaluation in AZ91 Alloy

Superplasticity determination in AZ91 sheets rolled at 400 °C with rolls not-heated and heated to100 °C, respectively, are given in Figures 3 and 4. It can also be seen that the lower rolls temperature, the higher elongation and lower flow stress. A profitable feature is that the as-rolled AZ91 sheets exhibit superplasticity at low temperature as 300 °C or 350 °C at strain rate of 0.0014 1/s, which is considered to result from grain refinement caused by the intense shear strain produced during DRSR process. Microstructural analysis is in progress which would provide detailed information for microstructure development during DRSR rolling. 140

400 AZ91 Alloy Rolls temperature= R.T. o Billet temperature= 400 C

120

AZ91 Alloy Rolls temperature= R.T. o Billet temperature= 400 C

(b)

100 Flow Stress, MPa

Total Elongation, %

300

(a)

200

80 60 40

100

20

Strain rate= 0.0014 1/s Strain rate= 0.014 1/s 0 275

300

325

350

375

400

425

0 275

450

Strain rate= 0.014 1/s Strain rate= 0.0014 1/s 300

325

350

375

400

425

450

o

o

Temperature, C

Temperature, C

Figure 3. Total elongation (a) and flow stress (b) vs. testing temperature in the AZ91 alloy rolled at 400 °C with the rolls kept at room temperature. 140

400 AZ91 Alloy o Rolls temperature= 100 C o Billet temperature= 400 C

AZ91 Alloy o Rolls temperature= 100 C o Billet temperature= 400 C

120

(b)

100 Flow Stress, MPa

Total Elongation, %

300

(a)

200

80 60 40

100

20

Strain rate= 0.0014 1/s Strain rate= 0.014 1/s 0 275

300

325

350

375

400

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450

0 275

Strain rate= 0.014 1/s Strain rate= 0.0014 1/s 300

325

o

Temperature, C

350

375

400

425

450

o

Temperature, C

Figure 4. Total elongation (a) and flow stress (b) vs. testing temperature in the AZ91 alloy rolled at 400 °C with the rolls heated at 100 °C.

3.4

Superplasticity in Evaluation in AZ31 Alloy

Figures 5 and 6 display the result of superplasticity evaluation in the AZ31 alloy sheets rolled at 300 °C with rolls temperature of room temperature and 100 °C, respectively. The effect of rolling and tension testing conditions on the total elongation and flow stress shows similar feature as that in AM60 alloy. But the AZ31 alloy deformed specimens are not uniform and with visible necking at the investigated strain rates and temperatures, in spite of the fact that some specimens exhibited large total elongation to fracture as more than 200%. The reason is still unclear at present time. Superplastic characteristics of the as-rolled AZ31 sheets are still under investigation in a wide strain rate and temperature range.

304 300

100 AZ31 Alloy Rolls temperature= R.T. o Billet temperature= 300 C

(a)

80 Flow Stress, MPa

Total Elongation, %

200

100

AZ31 Alloy Rolls temperature= R.T. o Billet temperature= 400 C

60

40

20

Strain rate= 0.014 1/s Strain rate= 0.0014 1/s

Strain rate= 0.0014 1/s Strain rate= 0.014 1/s 0 275

300

325

350

375

400

425

0 275

450

(b)

300

325

o

350

375

400

425

450

o

Temperature, C

Temperature, C

Figure 5. Total elongation (a) and flow stress (b) vs. testing temperature in the AZ31 alloy rolled at 300 °C with the rolls kept at room temperature. 300

120 (a)

100

200

Flow Stress, MPa

Total Elongation, %

AZ31 Alloy o Rolls temperature= 100 C o Billet temperature= 300 C

100

300

325

350

375

400

o

Temperature, C

425

450

(b)

80 60 40 20

Strain rate= 0.0014 1/s Strain rate= 0.014 1/s 0 275

AZ31 Alloy o Rolls temperature= 100 C o Billet temperature= 400 C

0 275

Strain rate= 0.014 1/s Strain rate= 0.0014 1/s 300

325

350

375

400

425

450

o

Temperature, C

Figure 6. Total elongation (a) and flow stress (b) vs. testing temperature in the AZ31 alloy rolled at 300 °C with the rolls heated at 100°C.

3.5

Superplastic Forming

Figure 7 shows a macrograph of domes blow formed from AM60 alloy sheet rolled under condition of 400 °C / R.T./ 5%. The forming test was conducted within the respective optimal superplastic deformation temperature range depicted by Figure 1, and it seems

Figure 7. Hemi-spherical AM60 alloy domes blow formed at 425 °C under the constant pressure of 0.68 MPa held for 2 minutes (left) and 6 minutes (right), respectively.

305 possible to shape DRSR-Mg sheet components by SPF technique. Up to present, the SPF magnesium research is limited, despite that very high elongation (>400%) at rather low temperature (<300 °C) has been demonstrated in wrought magnesium alloys [10]. But the materials used in these studies are not produced by conventional bulk working method which means costly for large volume production. The current superplastic forming test implies that it might be a suitable way to produce magnesium sheet by DRSR process for superplastic forming. Further investigation is useful to determinate more appropriate forming condition to reduce the production cost as low as possible.

4

Summary

AM60 alloy, AZ91 alloy and AZ31 alloy sheets with appropriate mechanical properties could be produced from cast billets by a different-roll-speeds-rolling (DRSR) process applied in this study. The intense shear strain introduced through the sheet thickness during DRSR operation is thought to induce grain refinement and consequent enhancement of superplasticity in the examined magnesium alloys. The as-rolled AM60 alloy sheet was successfully gas-blow formed under the superplastic deformation conditions illustrated by tension test, indicating that DRSR has the potential to manufacture magnesium alloy sheet for further superplastic forming. DRSR is an unexploited technique that would provide many subjects for investigation, systemic and detailed work is necessary either for generating scientific knowledge or for application of the process.

Acknowledgement One of the authors (S.L. Dong) greatly acknowledges the finical support of Japan Society for the Promotion of Science (JSPS).

References [1] N. Hort, K. U. Kainer in AIST-ISEM International Symposium 2003, Nagoya, Japan, p.17 [2] Int. Magnesium Association, http://www.intlmag.org [3] R.E (Bob) Brown in Magnesium Technology 2002 (Ed: H. I. Kaplan), TMS, 2002, p.155 [4] P. E. Krajewski in Magnesium Technology 2002 (Ed: H. I. Kaplan), TMS, 2002, p.175 [5] T. Sakai, S. Hamada, Y. Saito, Scripta mater. 2001, 44, 2569-2573 [6] S.-H. Kim, J. K. Lee, D. N. Lee in Ultrafine Grained Materials II (Ed: Y. T. Zhu, T. G. Langdon, R. S. Mishra, S. L. Semiatin, M. J. Saran, T. C. Lowe), TMS, 2002, p.55 [7] Q. Cui, K. Ohori, J. J. Metals. 2003, 73(3), 214-219 [8] J. S. Vetrano, JOM, 2001,3, 22 [9] R. Grimes, Mater. Sci. Tech. 2003, 19(3), 3-10 [10] H. Watanabe, T. Mukai, K. Higashi, in Ultrafine Grained Materials II (Ed: Y. T. Zhu, T. G. Langdon, R. S. Mishra, S. L. Semiatin, M. J. Saran, T. C. Lowe), TMS, 2002, p. 469

Results of Tempered Hydroforming of Magnesium Sheets with Hydroforming Fluids Reimund Neugebauer, Michael Seifert Fraunhofer Institute of Machine Tools and Forming Technology in Chemnitz, Germany (IWU)

1

Introduction

Lightweight design is becoming increasingly important in such areas as the automobile industry due to the necessity of keeping resource utilization to a minimum while reducing environmental pollution. But, this not only includes taking advantage of the benefits of condition and structural lightweight design. Lightweight materials are becoming increasingly important in high-strength and super high strength steels, aluminum and magnesium sheet metal. Unfortunately, this proven potential for weight-saving is frequently linked to limited deformation capacity as compared with conventional materials. This is the reason why tempered deformation has become increasingly important in the last few years. Preliminary tests (/1/.../2/) have indicated an increase in deformation capacity of 250% to 350% coupled with a reduction in the deformation force required of 40% – 50% with magnesium wrought alloys in a temperature ranging from 220° to 250 °C (/3/.../5/).

2

The Requirements Made of Tempered IHB of Mg Components

Pure research has indicated the following requirements for the application of tempered hydroforming sheet metals for machining blanks made of magnesium wrought alloys. We primarily focused our definition on the specific plant design, maximum controlled temperatures and process engineering for utilizing the maximum effects. Since we wished to have a quick heat transition from the deformation medium to the component without making external heating equipment necessary, we decided to eliminate gaseous deformation media. Our range of tests required temperatures ranging up to 330 °C and working pressures up to and including 800 bar. And in general, it is indispensable to have reproducible process parameters that can be adjusted precisely over the entire range of the study to produce a meaningful study of the deformation process. A final constraint is ensuring that the unit has to be UVV compatible.

3

The Tempered IHU/IHB Unit

Figure 1 illustrates the basic functioning and the structure of a tempered IHU/IHB unit that satisfies these requirements. This unit consists of an electrical heating unit with two separate cycles. One cycle is used for heating the lower and upper part of the tool and

307 heating the low-level service reservoir (in the form of an integrated heat exchanger). A thermal oil was used as the fluid with a temperature resistance ranging up to 355 °C. The fluid is suctioned by a special pump in the second (high-pressure) cycle and pressed through the special precision filter at 16 bar, filtered to 10 μm and then conducted to the secondary side of the tempered pressure intensifier where the oil is compressed to 800 bar and the component deformed at this pressure. A temperature regulator is built into the return line for comparing the actual temperature with the target. We chose a thermal oil for the high-pressure cycle with a temperature resistance ranging up to 330 °C and a high oxidation stability since the fluid comes into contact with oxygen from the air (for instance, when rinsing or opening the tool). The tempered hydroforming tool may be impinged with a different temperature than the deformation medium, which involves certain effects at higher temperatures as far as the softening capacity of the Mg material is concerned, in particular in the pressure pad range.

Figure 1: A tempered hydroforming unit (a very simplified schematic diagram)

4

The Tempered Hydroforming Tool

Figure 2 illustrates the design of the tempered hydroforming tool. the upper part of the tool

the forming engraving the tool box (the upper and lower part )

the lower part of the tool

heating ducts

the high-pressure medium forward and return run

Matrize stainless steel perforated plate (circumference) flexible heat insulation board

heat insulation board 1 heat insulation board 2

Figure 2: A cross-section of the tempered hydroforming-tool (the front view and top view) indicating the heating channels, the high-pressure feed and drain channels and thermal decoupling used

308 The process of thermal decoupling takes place between the internal smart components and the tool block on the one hand and between the tool block and the base plates on the other hand, thus ensuring that it remains within the permissible maximum temperature of 70 °C for the press table and the chisel.

5

Tempered Hydroforming Deformation Tests

In our tests on tempered hydroforming sheet metals with liquid active means, we studied the magnesium materials listed in Table 1 and the aluminum material AlMg4.5Mn0.4 (AA 5182) by way of comparison. However, our further comments only refer to the results of our work with the material AZ 31 B. Table 1: Commercial composition of AZ 31 B, MA2-1 and MA2-1 HP sheet metal Material-spectrum AZ 31 B lt. ASTM-Standart AZ 31 B (*) MA 2-1 (*) (**) MA 2-1 HP (*) (**)

Al 3,00 2,78 3,90 4,30

Zn 1,00 0,96 1,35 1,20

Mn 0,35 0,06 0,50 0,30

Fe <0,005 0,010 0,020 0,005

Cu <0,05 0,005 0,018 0,008

Ni <0,005 <0,001 0,002 0,001

(*) Readings (**) surface chromatised

We studied tempered hydroformability of blanks made of AZ 31 B using the interior component of a motor hood (on a scale of 1:4). There were 2-dimensional cracks at an internal pressure of 155 bar and an approximate drawing depth of 10 mm at room temperature (refer to Figure 3). This demonstrates that it is not possible to manufacture these components without tempering. crack in the corner radius area

quality components

constrictions , cracks in the corner radius area

crack

20 °C

150 °C

220 °C

280 °C

Figure 3: Internal component of the motor hood of a passenger car made of AZ 31 B (s0 = 1.3 mm), deformation results as f (t, material)

We got the best results (quality components) at deformation medium temperatures ranging between 180° – 220 °C and tool temperatures ranging between 180° – 200 °C (refer to Fig. 3 and Fig. 4). There were brittle cracks at low temperatures ranging between 20° 130 °C and significant cracks were detected in the plateau and corner radius area at temperatures ranging between 130° - 150 °C. Finally, smaller cracks could be observed in the inner circumferential corner radiuses at temperatures ranging between 150° - 180 °C. There were local constrictions in the inner corner radiuses at higher temperatures ranging between 220° - 250 °C that extended to small, albeit locally limited cracks at even higher temperatures (approximately 250° – 280 °C).

309 Momentary System Limits for the Tempered Hydroforming Unit (pi max.= 800 bar; tmax.= 280 °C)

900

inner pressure pi [bar]

800 700 600

cracks Reißer

large brittle cracks

500

significant cracks

400

small cracks in the corner radiuses

excellent component range

constricsmall tions in the cracks in corner the corner radiuses radiuses

insufficient component formation

300 200 100 0 20

40

60

80

100

120

140

160

180

200

220

240

260

280

300

deformation temperature [°C]

in. clamping force Fs min. = 17.000 kN

Figure 4: The process parameter window of the interior component of the motor hood (AZ 31 B, s0 = 1.3 mm, Rmin comp. = 5.0mm, drawing depth of 27 mm, major stretch forming stress: ϕ1 max = 0.23; ϕ2 max = 0.12)

6

Analyzing the Components

6.1

Varying Wall Thicknesses in Selected Cross-Sections and a Comparison with the Results of FEM Simulation

We measured varying wall thicknesses along selected cross-sections as can be seen in Figure 5. These cross-sections were selected to run through the critical area where there were folds (corner or intake areas) or cracks (the inner radiuses) in the test series. The maximum sheet metal thickening was 1.43 mm in the area directly next to the pressure pad zone. The high level of clamping force (as much as 40,000 kN) displaced the material in the form of a bead into this zone from the area of the pressure pad. There were the greatest reductions in sheet metal thicknesses of roughly 30% at the inner circumferential radius (items 205 and 460 mm). The inner areas of positions 210 – 450 mm exhibited an even reduction in sheet metal thickness of roughly 8%. FEM-Programm: PamStamp (s o =1,3 mm; t = 220°C)

Position des Schnittes C

A

[mm] Blechdicke [mm] thickness sheet

Change in Sheeting Thickness (Schnitt (Cross-Sections A – C) Blechdickenverlauf A-C) Werkstoff:AZ AZ 31 Material: 31B B

1,5 1,4 1,3 1,2 1,1 1 0,9 0,8 0,7 0,6 0,5 0,4 0,3 0,2 0,1 0 0

50

simulation

100

150

200

actual component

250

300

Position [mm] position [mm]

350

400

450

Simulation

500

550

Realteil

Figure 5: A comparison of the wall thickness variation in the interior component of the motor hood made of AZ 31 B along the A–C cross-section with the FEM simulation results.

310 6.2

Component Accuracy and Resilience Capacity

We used a 3-D coordinate measuring instrument to determine the differences between the targeted and actual radiuses and the differences in geometries as a mean of determining the component accuracy along defined cross-sections. Figure 6 shows the results in terms of the forming accuracy of the radiuses. Cross-sections B-B and C-C only differed by < 2%, while the two inner circumferential radiuses (measuring points A5 and A8) differ by as much as 18% (a tolerance of +/- 10%) in cross-section A-A (transverse to the rolling direction). Since the radiuses of the two measuring points along cross-sections B-B and C-C are within tolerances, we may postulate that there was resilience in the two areas of the radiuses mentioned above. If we neglect this difference due to twisting in the component when cooling down to room temperature, we may designate this as acceptable forming accuracy in the process of tempered hydroforming magnesium sheet metal. point

ideal value [mm]

cross-sections A-A actual value [mm]

variance [% ]

5,0 10,0 10,0 5,0 5,0 10,0 10,0 5,0 5,0 10,0 10,0 5,0

5,4 10,3 10,0 5,4 5,9 10,4 10,4 5,9 5,4 10,0 10,3 5,3

8,0 3,0 0,0 8,0 18,0 4,0 4,0 18,0 8,0 0,0 3,0 6,0

deviation < 5,0 %

deviation < 10,0 %

deviation > 10,0 %

A1 A2 A3 A4 A5 A6 A7 A8 A9 A10 A11 A12

C

B B1 B2 B3 A2

A3

A6

B4

A7

A10

A11

A

A A1

A5

A4

B5

A8

A9

A12

B6 B7 B8

C

B

position of the measuring points

Figure 6: Comparing the targeted radius to the actual radius at selected cross-sections

6.3

The Change in Mechanical Properties after Tempered Hydroforming

Something else of interest with tempered hydroforming is how temperature impingement modifies the mechanical properties during deformation and what properties the component has after it cools down to room temperature. This is especially important with components that might be involved in collisions as in this case. Therefore, we also determined the mechanical indicators of the initial magnesium sheet metal blank and the tempered hydroformed components at room temperature in a variety of ranges (refer to Figure 7). position of the measuring points

43,50 42,43 42,40

measurements made with RT

modulus of e lasticity [M Pa]

1 2 3

4,10 4,60 5,60

34,30 35,90 38,30

4

4,80

36,10

specimen 2

(across the RD)

tensile specimen: 180 x 20 mm

specimen

13,35 15,73 13,79

Ag [%]

specimen 3

0° 45° 90°

modulus of sample e lasticity numme r [M Pa]

(along the RD)

Ag [%]

specimen 1

rolling dire ction

4

te mpe re d hydroforme d motor hood

initial she e t

rolling direction

Figure 7: Comparing the mechanical indicators on the inside component of the tempered hydroforming motor hood at four different points with the initial blank (AZ 31 B)

311 The modulus of elasticity drops by an average of 28.9% (12.4 Mpa) due to tempered hydroformed deformation while the average remaining uniform elongation is 4.8% after the component has cooled off. The following simulations have to demonstrate whether the remaining elongation is sufficient for a possible collision. 6.4

The Surface Quality before and after Tempered Hydroformed Deformation

The object of these studies was to assess the quality of the exterior skin of the magnesium sheet metal, how well it can be lacquered and its tribological properties. As can be seen in Figure 8, the surface topography of the initial sheet metal is very level, smooth, even and has a slight roughness. initial sheet Rz 0° = 3,46 µm Rz 45°= 2,33 µm Rz 90°= 3,16 µm

rills

case sheet

Rz 0° = 3,95 µm Rz 45°= 4,19 µm Rz 90°= 3,15 µm

plateau range

middle range

Rz 0° = 4,15 µm Rz 45°= 3,21 µm Rz 90°= 3,00 µm

Rz 0° = 4,15 µm Rz 45°= 3,21 µm Rz 90°= 3,00 µm

scratches

Figure 8: 3-D surface roughness of the inside component of the tempered hydroforming motor hood (AZ 31 B) in various component areas

When we compare the 3-D surface roughness of the initial sheet metal with surfaces after tempered hydroformed deformation in a series of typical areas, we can observe a visible roughening of the surface in the case and plateau area due to the significant degree of deformation there. Likewise, we detected slight furrows and scratches. However, in the middle area, there is significantly less roughening, making it almost comparable with the initial surface. All in all, we can say that the surface remains relatively homogenous and we were not able to detect an „orange peel“ forming as some aluminum sheet metals at higher degrees of deformation. The Rz readings indicate that the surface was only roughened to a very slight degree due to the deformation process on an average of 1 μm and it was not possible to identify any influence on the rolling direction at the roughening point. Therefore, we may assess the surface quality and tribological properties of the AZ 31 B sheet metal as excellent and place them on the same level as AlMg4.5Mn0.4 sheet metal. 6.5

Studies of the Microstructure and REM Pictures

Grinding samples were made of the initial sheet metal and the deformed component in microstructure studies to determine the mean grain size. It was not possible to detect any significant differences in the mean granulation diameter or microstructure regarding the

312 rolling direction. Both are a fine-grained homogenous microstructure (with a mean granulation diameter of 10 μm) without any irregularities. Likewise, it was not possible to detect a deformation microstructure, impurities, occlusions or irregularities in the microstructure, etc. Figure 9 is image from the scanning electron microscope. This Mg alloy has a 3-layer system according to /6/ with at least two different phases. We can see darker second phases (a phase rich in Al or Zn), the proportion of which corresponds to the ternary phase diagram of the Mg alloy. There is a trend towards traces of the formation of microcracks that are of intercrystalline nature or that take place an inner phase border. This potential of microcrack formation may lead to preliminary constriction and component failure due to crack formation with the deformation of the component. The crystalline surface topography shows a fine-grained isotropic microstructure, although the tempered hydroforming process did not produce any anisotropies.

(second phase)

(large-angle granulation limits: danger of intercrystalline formation of microcracks )

(contrast differences )

(grain boundary) (planar sliding stage (bands) with varying orientation)

Figure 9: Scanning electron microscope images of the surface (10 μm) of the inside component of the tempered hydroforming motor hood in the area of the blank (AZ 31 B, S1 ≈ 1.21 mm, WD = 15 mm, EHT = 20.00 kV)

7

Reference

[1] Doege, E.; Umformen von Magnesiumwerkstoffen, DGM Fortbildungsseminar, Clausthal-Zellerfeld 1998. [2] Dröder, K.; Untersuchungen zum Umformen von Feinblechen aus Magnesium legierungen, Dissertation, IFUM Hannover, 1999. [3] Solberg, J. K. et al.; Superplasticity im magnesium alloy AZ 91, Materials Science and Engineering, A 134, 1991. [4] Böndel, B.; Magnesium im Automobilbau, Ingenieur-Werkstoffe, 4 (1992). [5] Kainer, K. U.; Potential von Magnesium in der Karosserie - Möglichkeiten und Entwicklungsbedarf, 5. Euroform - Bonn, Februar 1998. [6] Kammer, C. u.a.; Magnesiumtaschenbuch, Aluminium-Verlag Düsseldorf, 2000, 1. Auflage.

Bending Magnesium Extrusions R. Gradinger, M. Brandecker, H. Kilian, A. Wahlen ARC Leichtmetallkompetenzzentrum Ranshofen GmbH (LKR), A-5282 Ranshofen

Abstract Experiments of a research program concerning the bending of hollow extrusions of the magnesium wrought alloys are described. The poor formability of standard alloys like AZ31 at room temperatures is overcome by forming at elevated temperatures. The appearance of twinning is observed by optical micrographs and proved by pole figures.

1

Introduction

Magnesium alloys have aroused interest in the recent years from industry. Especially cast alloys – predominantly AZ91D – found wide applications in the automotive industry e.g. as an instrument panel beam or as a gear box. Recently some components of wrought Mg alloys for automotive applications were presented in the prototype of the Volkswagen 1liter-car and Porsche’s new Carrera GT. For an economic application of magnesium wrought alloys in large series production the whole process chain has to be available and the quality of manufacture has to be competitively priced. LKR treats the process chain of magnesium alloys in the fields alloy development, vertical continuous casting, component casting, rolling, extrusion, forming, joining, heat treatment and recycling [1-4]. Analogously to sheet drawing [4] large plastic deformations of extrusions by forming can not be achieved with standard procedures at room temperature derived from processing aluminium alloys or steel. One possibility to increase the ductility of magnesium alloys is to increase the forming temperature above 200 °C.

2

Formability of magnesium alloys

Caused by their hexagonal atomic lattice the formability of magnesium and most of its alloys is limited by the low number of slip planes available at room temperature. Dislocations are hindered in their movement and twinning accounts for a major amount of the deformation. An improvement in ductility can be achieved when the forming is performed at temperatures above 200 °C because additional slip planes can then be activated [5]. Another way of achieving increased ductility is to drastically reduce the average grain size. Then the deformation is dominated by slipping along the grain boundaries. In [6] Emley presented some results (after Chapman and Wilson) which are shown in Figure 1: the transition from low to high formability (given as reduction in area) moves to lower temperatures as the grain size decreases.

314

Figure 1. Influence of grain size on the formability [6].

3

Figure 2. Yield strength Rp0,2, ultimate tensile strength Rm and fracture elongation A versus temperature of AZ31 sheet metal [4].

Bending of Extrusions

Commercial available magnesium extrusions out of AZ31 typically have a grain size between 10 μm and 200 μm with medium grain sizes of about 30 μm to 80 μm. According to Figure 1 the reduction of area will be much lower than 20% at room temperature. This is in good accordance with results from tensile tests performed with 2.0 mm AZ31 sheets [Figure 2]. The main aim of the bending trials with hollow extrusions was to produce a 90° angle with radii of 200, 300 and 500 mm, respectively. A second aspect was the surface quality after bending. The samples were hollow extrusions with a cross section of 40 mm x 40 mm and a wall thickness of 3 mm. After filling the hollow section with quarry sand it was compacted and closed at the ends with matching aluminium parts. The 3-roll bending machine in horizontal design was chosen to perform these tasks. The main advantage of this type of device is the low costs for adapting to different extrusion geometries. Nevertheless alternative processes like die bending or producing bent work pieces by combined extrusion and bending [7] are under consideration. The used Arkus 12 type bending machine possesses measuring as well as programming capabilities and calculates the optimal single bending steps to fulfil the required parameters. Directly before bending the filled profiles as well as the tools of the bending were heated up to varying temperatures. The temperatures of the profiles as well as the rolls were measured at the beginning and the end of the bending process – in case of the profiles at three locations at both ends and the centre of the sample, respectively. Table 1 lists process parameters and results of one experimental run – beginning with 260 °C down to 180 °C starting temperature of the filled extrusion. The hydraulic pressure of the bending device was recorded at the beginning and the end of the bending procedure, the difference between these to values is given in the column “Press(ure) rise”. The adjusted bending radius of 300 mm over 90° angles was not achieved exactly – the missing compliance is due to uncertainties in the elastic spring back and the shortening in the compression zone by twinning. An important parameter is the temperature drop during the duration of the experiments: due to the stepwise bending process the profiles lost between 55 and 80 °C while the heated rolls cooled down from 230 to 80 °C. In inverse proportion

315 the system pressure increased from 15 to 30 bar comparing the first temperature level with the last one. The pressure increase within a single test is due to hardening with increasing plastic deformation; the bending at lower temperatures results in a higher pressure necessary for the plastic deformation in the bending zone. Analogous results were found for R = 500 mm and R = 200 mm which are not presented here. In order to estimate the amount of twinning in the whole plastic deformation pole figures were measured for two planes in the crystal lattice: (1000) and (0002). Regarding the base material the direction of the c-axis of the single grains is parallel or 90° perpendicular to the extrusion direction with a scatter of only about 10°. Comparing the pole figures of samples formed at room temperature to samples as fabricated a significant amount of grains changed its orientation by twinning in the compression region. This effect does not appear in this amount for the bending process at a target temperature of 260 °C. Theoretically no twinning should occur at temperatures above 200 °C, but since the temperature reached 200 °C at the end of the process some twinning contributed to the forming. Table 1. Parameter and results of the bending tests (radius 300 mm, angle 90°)

Target temp. [°C]

Actual temp. top/centre/end [°C] @ begin

Actual temp. top/centre/end [°C] @ end

Temp. drop top/centre/end [°C]

Tool temp. Hydr. begin-end Press. [bar] [°C]

Press. rise [bar]

Bent radius [mm]

260 240 220 200 180

300/260/300 260/240/260 240/220/240 220/200/220 195/175/195

240/200/240 200/160/200 180/150/180 160/130/160 140/115/140

60/60/60 60/80/60 60/70/60 60/70/60 55/60/55

230/140 140/130 120/110 100/90 90/80

0/5 2/8 2/10 2/13 2/13

293 295 295 297 300

15/20 17/25 17/27 17/30 17/30

Table 2. Pole figures of different material conditions

Base material (as fabricated)

(1000)

(0002)

Compression region Tension region Forming @ RT Forming @ RT

Compression region Forming @ 260°C

316 These results support the micrographs shown in Figure 3 to Figure : after bending at room temperature or bending at elevated temperatures not higher than 200 °C (the minimum temperature of the sample in Figure was 160 °C, the one of Figure 200 °C) twinning can easily be found in the compression zone. In order to avoid twinning at all forming temperatures below 200 °C have to be avoided.

Figure 3. Micrograph in extrusion direction; sample not deformed.

Figure 4. Micrograph perpendicular to extrusion direction; sample not deformed.

Figure 5. Micrograph perpendicular to extrusion direction; deformed at room temperature; tension region.

Figure 6. Micrograph in extrusion direction; sample deformed at RT; compression region.

Figure 7. Microgr. in extrusion direction; sample deformed at 260 °C; compr. region.

Figure 8. Micrograph in extrusion direction; sample deformed at 240 °C; compre. region.

317

4

Applications

The results of these experiments were used to get more insight into the bending of Mg extrusions, so some other profiles used in automotive prototypes or automotive component assemblies were as well bent. One of these extrusions is shown in Figure 3. This magnesium profile is a double chamber hollow extrusion described in more detail in [8]. Clearly visible are some markings of the rolls which can be avoided by more detailed adoption of the roll’s geometry.

Figure 3. As extruded and bent profiles for a magnesium car bumper.

5

Conclusion

Hollow extrusions of standard magnesium alloys can be bent to conventional shapes by 3roll bending by applying the deformation under elevated temperatures. During the forming operation the temperature should not reach 200 °C or less. By adjusting and cleaning the rolls severe markings can be avoided. Further improvements are subject to R&D work on this topic.

6

References

[1] F. Pravdic, H. Kilian, D. Leitlmeier, A. Wahlen, P. J. Uggowitzer, B. Vogler, P. Wohlmuth, Wire 2003, 1, 37 –41. [2] R. Potzinger, H. Kaufmann, P. J. Uggowitzer in Proc. 6th Int. Conf. on Semi-solid Processing of Alloys and Composites, Turin, 2000, 85-90. [3] U. Galovsky, M. Kühlein, in Magnesium Technology 2001 (Ed.: J. Hryn), The Minerals, Metals & Materials Society, Warrendale, PA, 2001, 49-53. [4] R. Gradinger, P. Stolfig in Magnesium Technology 2003 (Ed.: H. Kaplan), The Minerals, Metals & Materials Society, Warrendale, PA, 2003, 231-236. [5] C. Kammer et al., Magnesium-Taschenbuch, Aluminium-Verlag, Düsseldorf, Germany, 2000, 241. [6] E. F. Emley, Principles of Magnesium Technology, Pergamon Press, Oxford, Great Britain, 1966, 490. [7] M. Kleiner, EP 0 706 843 B1, 1999 [8] K. Hausberger, B. Götzinger, T. Friesenbichler, Mobiles, 2002, 28, 58-63.

Drawing of Magnesium Wires at Ambient Temperature Jörg Eickemeyer, Albert Güth, Martina Falter, Ralph Opitz Leibniz-Institut für Festkörper- und Werkstoffforschung Dresden, Institut für Metallische Werkstoffe

1

Introduction

Hexagonal close packed (hcp) metals like magnesium or cobalt behave little ductile under conditions of cold forming at ambient temperatures. This is attributed to the very restricted number of independent slip systems which can be activated at room temperature [1]. Basal slip is the predominating mechanism of plasticity and pyramidal or prismatic slip should become increasingly important only at higher temperatures [2]. Therefore, bulk forming of magnesium is generally performed at temperatures of at least about 200 °C [3]. This holds for deformation processes which work under the action of high pressures, like extrusion, forging or rolling. On the other hand, when tensile stresses are applied during magnesium working it behaves much more sensitive with respect to cracking phenomena. Hence, wire drawing of magnesium and its alloys as a cold forming process in which tensile stresses are necessarily included needs very special knowledge. Especially, the material flow very close to the surface of the work piece seems to be substantially important, as recently shown [4]. This experimental investigation was aimed at a better understanding of the near surface material flow and the convenient tribological conditions for a successful cold drawing process in case of pure magnesium metal.

2

Experimental

Experiments were performed with magnesium of 99.9 mass-% purity. The initial rod material was extruded at 350 °C from about 23 mm in diameter to 3.85 mm. Th extruded wire of 3.85 mm diameter was directly used in the drawing tests. The initial surface roughness was Ra = 0.7 μm for the as extruded state. Further, magnesium was hot extruded from 25.4 mm to 6 mm in diameter. Subsequently it was mechanically peeled and specimens of 5.70 mm in diameter were prepared (Ra ≈ 0.1 μm). Cold drawing of rods and wires, respectively, was carried out under laboratory conditions. It proceeded on a drawing bench at a drawing rate of 1...6 m/min. The specimens were maximum two metres long. Cemented carbide dies with an opening angle of 2α = 12° or 2α = 8° were used without die cooling. The reduction of area per draft was in the range from ε = 3.6 % to ε = 26%, as obvious from Table 1.

319 Table 1. Different reductions of area and tool angles used in magnesium wire drawing.

Sample

Wire diameter /reduction of area

Drawing angle 2α

1

∅3.85  3.54 mm / ε = 15.5%

2α = 8°

2

∅3.85  3.54 mm / ε = 15.5%

2α = 12°

3

∅3.85  3.72 mm / ε = 6.6%

2α = 12°

4

∅3.85  3.72 mm / ε = 6.6%

2α = 8°

5

∅3.85  3.38 mm / ε = 22.9%

2α = 12°

6

∅3.85  3.38 mm / ε = 22.9%

2α = 8°

The surface topography of the forming good was investigated by scanning electron microscopy (SEM) and roughness measurements. The roughness was measured by means of an laser automatic focus system as well as by using a mechanical stylus. Roughness was regarded to all wave lengths smaller than 250 μm. The peak roughness parameter Rz(DIN) was used to characterise the roughening and cracking processes, respectively. The mainly used lubricant was basically a lime soap dispersion in water, which was applied as a solid lubricant (Ko2) and partially also as a liquid lubricant (Ko3) [5]. The dispersion was converted to a solid lubricant film by means of dipping the wire to be drawn into the hot aqueous dispersion and drying it subsequently at increased temperatures. The main components of this lubricating system were calcium stearate, a non-ionic detergent and a layer forming substance based on a natural product. The liquid films were used for thin film lubrication and the dried ones for thick film lubrication. In addition, some mineral oil lubricants were tested in magnesium wire drawing, too.

3

Results and Discussion

Application of the liquid lubricant dispersion Ko3 and a mineral oil lubricant, respectively, was not successful in the drawing experiments. The lubricant film broke down and led to a cold welding between the wire and the drawing tool resulting in die pickup and longitudinal wear tracks on the wire surface. This tribological behaviour of magnesium was absolutely different from that one in cobalt drawing, where thin film lubrication enabled the wire drawing process [6]. However, the application of a coating with solid soap Ko2 led to a stable lubrication regime. Starting from ∅ 5.70 mm the initially peeled magnesium wire could be drawn to ∅ 5.20 mm, ∅ 4.90 mm, ∅ 4.30 mm and ∅ 3.90 mm. The total reduction of area was εg = 53.2% and single reductions of area were up to εe = 23.0%. As obvious from Fig. 1 wear tracks were not formed. However, as expected, there was also a tendency to shear localisation and shear crack formation. But this tendency was clearly less pronounced than in cobalt drawing. The over 4 steps drawn magnesium wires formed cracks in the surface predominantly under 30° to the drawing direction (Fig. 1). Despite this bad surface state an additional drawing step from ∅ 3.90 mm to ∅ 3.35 mm (εe = 26.2%) was performed without rupture or lubricant film break down. Consequently, the total reduction without intermediate heat treatment amounted to εg = 65.5%.

320

Figure 1. SEM micrograph of an initially peeled magnesium wire, drawn over 4 drafts from ∅ 5.7 mm to ∅ 3.9 mm (εg = 53.2%) using the solid lubricant Ko2.

The roughness measurements have shown that magnesium behaves less susceptible against surface irregularities than cobalt. Magnesium wires having an initial roughness peak parameter Rz(DIN) up to about 8 μm could be drawn without problems, i.e. magnesium could also be drawn in the as extruded state. Therefore, wire sticks were produced from the extruded wires and subsequently investigated in order to analyse the roughening processes along the whole deformed region. As apparent from Tabs. 2 and 3 the extruded initial wire surfaces had Rz values of 3.1 μm and 3.4 μm, respectively. Table 2. Roughness parameters Ra, Rz, Rt and Rmax of the magnesium wire stick 2; drawn from ∅ 3.85 mm to ∅ 3.54 mm (ε = 15,5%) using the solid soap lubricant Ko2; 2α = 12°.

Wire stick zone

Roughness parameters [μm] Ra

Rz

Rt

Rmax

Initial wire surface

0.7

3.1

4.8

3.4

Drawing cone region

1.0

4.8

7.4

7.0

Final wire surface

1.2

6.3

10.9

10.3

The unavoidable roughening during deformation was due to specific material flow in the near surface region and resulted in Rz = 6.3 μm for 2α = 12° and Rz = 7.8 μm for 2α = 8 μm (Tabs. 2 and 3). In all cases roughness developed as to expect [7]. The lower the tool pressure the more pronounced was the roughening. At one and the same reduction of area the lower drawing angle yielded the higher roughness parameters. Correspondingly, at one and the same drawing angle roughness values increased with increasing reduction of area.

321 Table 3. Roughness parameters Ra, Rz, Rt and Rmax of the magnesium wire stick 1; drawn from ∅ 3,85 mm to ∅ 3,54 mm (ε = 15,5%) using the solid soap lubricant Ko2; 2α = 8°.

Wire stick zone

Roughness parameters [μm] Ra

Rz

0,8

3,4

4,7

4,5

Drawing cone region

1,4

6,8

10,7

10,1

Final wire surface

1,5

7,8

12,6

11,9

Initial wire surface

Rt

Rmax

By the measurement of the surface topography it became evident that additionally to the general roughening process a surface waviness develops, as illustrated in Fig. 2.

(a)

(b) Figure 2. Roughness topography of the magnesium wire stick 1, drawn from ∅ 3,85 mm to ∅ 3,54 mm (ε = 15,5%) using solid lubricant Ko2; drawing angle 2α = 8°; (a) transversal and (b) longitudinal measurement.

322

Roughness parameter Rt [µm]

Whereas the roughening is clearly obvious from the transverse measurement (Fig. 2a) the longitudinal measurement shows more impressive the formation of surface waves (Fig. 2b). Before the wire enters the cone region there was no indication of any waviness. It is increasingly developed inside the deformation zone with increasing reduction of sectional area. The formation of surface waves in the frame of the primary profile was also the more pronounced the lower the tool pressure. Therefore, the tool with 2α = 8° was more convenient to demonstrate this phenomenon than the tool with 2α = 12°. It seems very likely that such a wavy surface morphology should substantially influence the stress and strain conditions during material flow, at least in the near surface region. The consequences for the crack development and drawability have to be investigated in more detail.

50

Initial wire surfac e

Deformation cone

Final wire surfac e

40 30 20 10

Wire stic k length Figure 3. Roughness development during a wire drawing operation for steel and magnesium.

The highly specific situation in magnesium wire drawing becomes more clear, if compared with steel wire drawing. As shown in Fig. 3 steel wire of some millimetres in diameter is drawn by means of lubricant carriers and soap powder starting with rather rough initial surfaces. During drawing the surface is smoothened. Comparable rough initial surfaces like for steel do not allow to draw magnesium. Magnesium can be drawn only starting from very smooth initial surfaces and it becomes rougher during drawing. The degree of roughening will be strongly dependent on the actual tribological regime. The control of the latter one has to be regarded as the crucial point in order to be successful in the drawing operation, at Rz > 12 μm the process failed [8].

323

4

Conclusion

The drawability of magnesium wire at ambient temperatures strongly depends on the roughness of the wire. The parameter Rz was used to characterise the roughness. At values of Rz > 12 μm magnesium was not drawable. Measurements of the roughness topography of magnesium wire sticks gave evidence that a waviness in generated inside the drawing whole. The consequences of this phenomenon for material flow and damage behaviour have to be investigated in more detail.

5

References

[1] Hayden, H. W.; Moffatt, W. G.; Wulff, J., The Structure and Properties of Materials – Vol. III, John Wiley & Sons, New York, 1965, S. 97-120. [2] Couret, A.; Caillard, G., Phil. Mag. A 63 (1991) 5, 1045-1057. [3] Doege, E.; Janssen, S.; Wiesner, J., Mat.-wiss. U. Werkstofftech. 32 (2001) 5, 48-51. [4] Eickemeyer, J.; Falter, M.; Opitz, R.; Güth, A., Int. J. Appl. Mech. Eng. 7 (2002) 3, 771-782. [5] Eickemeyer, J.; Reichert, J.; Vogel, H.-R.; Opitz, R, in: Applications and Materials Performance, Ed.: F. N. Smith et al., Canadian Institute of Mining, Metallurgy and Petroleum, Montreal, Quebec, Canada, 1997, S. 209-217. [6] Eickemeyer, J.; Falter, M.; Opitz, R.; Güth, A., Int. J. Appl. Mech. 7 (2002) 3, 771782. [7] Hill, R.; Tupper, S. J., J. Iron Steel Inst. (1948) Aug., 353-359. [8] Eickemeyer, J.; Falter, M.; Güth, A.; Opitz, R.; Reichert, J., Patent Application DE 101 42 204.0, 24.08.2001. This investigation was financially supported in the frame of DFG-Schwerpunktprogramm SPP 1074 „Erweiterung der Formgebungsgrenzen bei Umformprozessen“.

The MAGNEXTRUSCO project: Process and Alloy Development for Hydrostatic Extrusion of Magnesium A.J. den Bakker*, W.H. Sillekens*, J. Bohlen**, K.U. Kainer**, G. Barton*** * TNO Industrial Technology, Eindhoven ** GKSS Research Center, Geesthacht *** RWTH Aachen – Institute of Metal Forming (IBF), Aachen

1

Introduction

Magnesium is in the spotlight as a promising “new” industrial material. In the automotive sector, magnesium die-castings are already a recognized option because of their lightweighting potential along with their specific shaping possibilities and ease of manufacturing. As a further step, the incorporation of wrought magnesium parts is anticipated because of the enhanced material properties and the additional variety of products that can be made through forming processes such as extrusion. For extruded products, however, the use of magnesium is hampered by the current technical and economical limitations of the direct extrusion process. Within the European Community project entitled “Hydrostatic Extrusion Process for Efficient Production of Magnesium Structural Components” (acronym: MAGNEXTRUSCO), a consortium of industrial and research organizations is conducting research and technological development work to overcome the major obstacles in the production of magnesium extrusions. A major innovation in this respect is a new manufacturing route for magnesium sections, based on hot hydrostatic extrusion. This paper highlights a few results. The following section focuses on hydrostatic extrusion technology for magnesium, with the emphasis being on extrusion speed and surface quality. Next, some work on new wrought magnesium alloys with improved characteristics is presented. Additional information on the project is accessible through its homepage [1]. Also, a further contribution to this conference that gives a general introduction to the project is included in the symposium on research programs [2].

2

Process Development

In the hydrostatic extrusion process, a billet is forced through a die opening by pressurized fluid (the pressure medium). Unlike direct and indirect extrusion, there is no metallic contact between ram/container and billet. Features of the process are that plastic deformation takes place under high hydrostatic pressure and that there is almost no friction. As an example, Figure 1 shows how a single-hollow section is produced. This involves the forming of the work-piece around a mandrel, using a billet with a bore.

325 MAGNEXTRUSCO is centered round the development of the hydrostatic extrusion process for magnesium. Background to this is that current productivity of conventional magnesium extrusion is quite low. Actually, extrusion exit speeds of up to 20–30 m/min are reported for the easy extrudable alloys and simple shapes, but these values drop by a factor ten and more for the more demanding applications. A major limitation is the occurrence of hot shortness (or hot cracking). This surface defect is initiated by incipient melting of the material due to excessive temperature rise, which in turn results from the internal dissipation of mechanical work. As the required mechanical work in hydrostatic extrusion is comparatively low, the process window is larger than for conventional extrusion. The main objective for the hydrostatic extrusion of magnesium is thus set at reaching exit speeds of 50–100 m/min. Here, surface quality is an obvious consideration. As the current experience with hydrostatic extrusion is restricted to basic symmetrical shapes, extension to Figure 1. Hydrostatic extrusion of hollow shapes more complicated sections is another development item. For these investigations, the ASEA Quintus press at Boliden HME in Waalwijk (Netherlands) isused. This industrial press can provide a maximum fluid pressure of 1,400 MPa and has a container with an inner diameter of 0.19 m and length of 1.6 m. Initial extrusion trials were conducted with downsized billets (diameter 75 mm, length 500 mm). Although extrusion of these billets proved to be relatively easy, control of the process was difficult due to the very short cycle time. Additionally, the instantaneous release of stored energy in the pressurized fluid resulted in very high initial exit speeds. In some instances, high inertial force caused fracture of the extrusion. Improved control of the process was obtained with the extrusion of full-size billets (diameter 160 mm, length 790 mm). Peak pressures at the start of the extrusion cycle were lower or even absent and the increased cycle time left more scope for control during extrusion. 2.1

Extrusion Speed

As a means of benchmarking, hot hydrostatic extrusion trials were conducted with the commercial alloys M1, ZM21, and AZ31 [3]. The selected test profiles were a rod (∅ 15 mm), a tube (∅ 35 × 2.5 mm), an angle section (20 × 20 × 2 mm), and a so-called dumb-bell section (2 ∅ 6.5 + 25 × 3 mm). All extrusion trials were done at the same setting of the hydraulic system and thus of the ram speed. Downsized billets were used. The main extrusion data are summarized in Table 1. Basically, the extrusion ratio, the die angle, the billet temperature and the extrusion (exit) speed are process parameters that result in a certain extrusion pressure. The applied billet temperatures were lower than in direct extrusion, where 300–400 °C is customary. At these particular settings, no hot

326 shortness was observed from the surfaces of the extruded profiles, thus confirming that hydrostatic extrusion is suitable for processing magnesium at high speeds. Table 1. Experimental data for hydrostatic extrusion of magnesium test profiles (R: extrusion ratio; αD: die semicone angle; Tb: set-point billet temperature; v: nominal extrusion speed; p: steady-state extrusion pressure)

αD [°]

Tb [°C]

v p [m/min] [MPa]

490 590 720

TUBE: 12 55 12 55 12 55

285 220 285

25 25 25

590 850 850

720 780

DUMB-BELL: 30 30 290 30 30 290

62 62

620 690

αD [°]

Tb [°C]

v p R [m/min] [MPa] [–]

M1 ZM21 AZ31

ROD: 24 40 24 40 24 40

285 280 260

50 50 50

ZM21 AZ31

ANGLE: 55 30 55 30

280 285

114 114

ALLOY

2.2

R [–]

Surface Quality

Surface condition of the extrusions is an important quality aspect, as it affects postextrusion operations (forming, coating) and functional performance. Table 2. Surface imperfections observed in hydrostatic extrusion of magnesium

DEFECT

APPEARANCE

COUNTERMEASURES

Hot shortness (or hot cracking): cracks perpendicular to the extrusion direction due to incipient melting of the material

Reduce billet temperature and/or extrusion speed, so as to achieve a lower temperature of the material at the die exit

Scoring (or die lines): scratches longitudinally to the extrusion direction due to scoring in the cone and die

Alter tooling geometry and/or finish to improve metal flow through and reduce friction in the die

Roughening: irregular, coarse topography, occasionally accompanied by breaking out of particles

Use billets with more homogeneous and finer microstructure

Contamination: remnants of the pressure medium, difficult to remove when burned onto the surface or appearing on the inside of hollow shapes

Ensure correct sealing of press channel by billet and/or avoid complete extrusion of the billet

327

Extrusion speed v [m/min]

In the course of the investigations, some distinct surface imperfections were encountered for the various magnesium sections. Table 2 summarizes these defects, including the measures that may be taken to avoid them. Where hot shortness and scoring are the more generic surface defects in extrusion, contamination with the pressure medium – which is not even an actual surface defect, but evidently a problem – is typical for hydrostatic extrusion. Further, the surface quality of the M1 extrusions was generally found to be poor due to roughening. The billets of this alloy showed a very pronounced and coarse casting morphology (average grain size of 1–4 mm). It is assumed that the contrasting orientation of the grains with the flow pattern during extrusion has led to the severe surface problems with this alloy. As opposed to this, ZM21 and AZ31, with a much finer billet microstructure, were not affected by this particular defect. ZM21 Tube ∅35×2.5 mm Extrusion ratio R=74

138

External surface

35 μm

69

50 μm

45 μm

250 315 Billet temperature Tb [°C] Figure 2. Surface appearance and microstructure of magnesium, extruded at different parameter settings (longitudinal cuts)

As a further illustration of hot shortness, Figure 2 gives some optical micrographs of tubes produced from full-size billets with different billet temperatures and extrusion speeds. The surfaces, located left in each photo, correspond with the outside of the tubes. Inserted values denote the average grain size. Although not that severe as for the analogous example in Table 2, surface cracks appear for the high billet temperature and for the high extrusion speed. Here, the thermal loading in combination with the (frictional) stresses at the die exit leads to surface degradation. The conditions for which this occurs do obviously depend on alloy-specific properties such as flow stress and incipient (or nonequilibrium) melting temperature. These observations on hot shortness are backed up with finite-element simulations, using a model that includes the thermo-physical effects during hydrostatic extrusion. This involves the internal heat generation due to conversion of mechanical work on the one and the local heat exchange due to thermal conduction (within the work-piece and with the pressure medium and tooling) on the other hand. Further, a flow stress that depends on

328 strain, strain rate and temperature is incorporated. The model is implemented in the preand post-processor PEP and finite-element code LARSTRAN/SHAPE [4]. As an illustration, Figure 3 visualizes the calculated temperature distribution over the cross-section of the work-piece for extrusion of a tube after reaching steady state. Extrusion direction is from bottom to top, with the centerline at the left. Process parameters match those for the upper-left photo in Figure 2. The simulation predicts that a maximum temperature of 485 °C is reached at the external surface of the tube upon leaving the die exit. Analogously, a surface temperature of 475 °C is calculated for the lower-right photo, whereas the situation in the lower-left photo corresponds with a maximum temperature of 445 °C. For all three cases, the calculated internal surface temperature of the tube is 12 °C lower than the external temperature. These calculations thus demonstrate that the occurrence of surface cracking is associated with those processing conditions in which the local temperature reaches the higher values.

3

ZM21 Tube ∅35×2.5 mm Extrusion ratio R=74

Figure 3. Calculated temperature distribution in hydrostatic extrusion of magnesium (half crosssection)

Alloy Development

Within MAGNEXTRUSCO, alloy and extrusion process development are treated in close concert with each other, as there is an intricate interrelation between these. For one, alloy composition will influence feedstock quality and the behavior of the material during extrusion and through this will have a bearing on process performance as well as on the obtained product. Conversely, the extrusion process is basically a thermo-mechanical treatment that adds to the micro-structural evolution of the material and thus to the product properties. 3.1

Feedstock Quality

In addition to the investigation of commercial alloys like AZ31, emphasis is laid on the selection of new and modified alloys to improve feedstock quality. It is a concern that additions to these alloys are inexpensive as to enable later industrial use for bulk applications. Further, it is important that the cast-ability of these alloys allows for an industrial production of billets, utilizing the DC casting technique. The left-hand side of Figure 4 shows a typical cross-section of an AZ31 billet in as-cast condition. As compared to this, the billet microstructure of the new alloy on the right-hand side is much finer, which is noticeable also in the average grain size (inserts). Accordingly, more homogeneous extrudability and profile properties can be expected.

329 New alloy

AZ31

400–600 μm

100–200 μm

Figure 4. Microstructures of magnesium billets

3.2

Material Properties

Extruded rods were used for microstructural and mechanical evaluation. Figure 5 gives optical micrographs of some new alloys, along with their average grain size. The obtained microstructures are equi-axed and fine-grained.

6 μm

5 μm

Figure 5. Microstructures of extruded new magnesium alloys (longitudinal cuts)

Table 3 lists the as-extruded mechanical properties for the commercial alloys used in the trials and the newly developed ones (the latter being arbitrarily labeled #1, #2, and #3). Samples were tested in tension and in compression. As compared to M1, ZM21, and AZ31, the new alloys show an improvement in yield and strength values, and especially in ductility (elongation). Moreover, the anisotropy as expressed by the differences in yield stress between tension and compression is less pronounced in the new alloys, which is clearly an advantage.

330 Table 3. Mechanical properties of magnesium extrusions (TYS: tensile yield stress; UTS: ultimate tensile strength; CYS: compressive yield stress; UCS: ultimate compressive strength)

ALLOY

TENSION COMPRESSION TYS [MPa] UTS [MPa] Elongation [%] CYS [MPa] UCS [MPa]

M1 ZM21 AZ31 #1 #2 #3

192 175 198 192 235 207

268 258 278 237 273 302

12 23 23 32 25 25

86 116 155 171 164 174

396 418 418 364 388 414

For rod; longitudinal direction

4

Summary

MAGNEXTRUSCO is a research and technological development project that aims at the development of the hydrostatic extrusion process for efficient production of magnesium structural components. A consortium of industrial and research organizations from across Europe is working together in this field. Some of the major issues are alloy and extrusion process development; these are treated in close concert with each other. Finite-element simulations support the investigations. Trials with commercial alloys have demonstrated that extrusion speeds of 100 m/min and more can be achieved, depending on the alloy, shape and extrusion parameters. Surface quality is an important consideration in this respect. The newly developed alloys exhibit excellent strength and ductility, as well as a reduced anisotropy in mechanical strength. Acknowledgements MAGNEXTRUSCO is being conducted as a GROWTH project within the Fifth Framework Program of the European Commission under contract number G5RD-CT-200000424; the financial support is gratefully acknowledged. Further, the authors would like to express their appreciation to Boliden HME for their contribution and to the partners for their consent for publication.

5

References

[1] http://magnextrusco.rtdproject.net. [2] W.H. Sillekens, J. Bohlen, Proceedings of the 6th International Conference on Magnesium Alloys and Their Applications 2003 (symposium M: Research Programs). [3] W.H. Sillekens, J.A.F.M. Schade van Westrum, A.J. den Bakker, P.-J. Vet, Materials Science Forum 2003, Vols. 426–432 (Proceedings of THERMEC’2003), 629–636. [4] R. Kopp, G. Barton, Proceedings of EUROMAT 2003 (in press).

Microstructure and Mechanical Properties of Squeeze Cast AM60 Mg Composites Reinforced with AISI 304 Wires Jae Joong Kim, Min Soo Yoo*, Jong Dae Lim, and Nack J. Kim* Hyundai Motors Company, Whasung, Korea, *Pohang University of Science and Technology (POSTECH), Pohang, Korea

Introduction Metal/metal composite, one of the methods to overcome ceramic/metal composite in terms of elongation and strength, has been investigated in using metallic fibers as reinforcement (e.g., Al matrix composites reinforced with Fe alloy wires) [1-4]. Metal/metal composite has been expected to have better interfacial bonding strength than ceramic/metal composites. In addition, little difference in thermal expansion coefficient between reinforcement and matrix has been also expected to decrease tensile residual stress around the interface while exposing the composite to thermal cycles. Furthermore, metallic reinforcement can be expected to absorb the stress transferred from matrix without being easily fractured. However, intermediate phases, the reaction products at the interface between reinforcement and matrix, make the composite locally brittle and decohesive. Recently, Boland et al. have manufactured Al matrix composite reinforced with Ni alloy fibers by squeeze casting [5]. In the metal/metal composite, oxidized fibers are enabled to prevent the reaction between Al and Ni. Such composite has better tensile properties than the composite reinforced with no oxidized fibers. Additionally, Sun et al. show that the tensile stress in steel/steel composites manufactured by rolling process seems to obey rule of mixture (ROM) [6]. By recourse to the experimental results, reactivity and interfacial strength of matrix/reinforcement as well as mechanical properties of matrix and reinforcement make an important effect on the mechanical properties of metal matrix composite. The present study is aimed at developing Mg matrix composite with high strength/ elongation, and showing mechanical properties of metal/metal composites variable with those of reinforcement under fixing matrix and interfacial condition.

Experimental Procedures AM60 (Mg-6Al-Mn) alloy, one of commercial Mg alloys, has been selected as the matrix alloy since it has better ductility than AZ91 (Mg-9Al-1Zn) alloy. Reinforcement should have high strength, high melting point, elevated temperature stability and no intermediate phase with Mg (e.g., Ti alloy and Fe alloy) [7]. In this work, AISI 304 stainless steel has been selected as reinforcement since it has the above-mentioned characteristics. As-received AISI 304 wire with diameter 0.4 mm, produced by continuous drawing machine, has high strength (1.8 GPa), low elongation (1.0 %) and clean surface with the removed scale by surface treatment. The fabrication process for AISI 304 reinforced AM60 Mg composite is squeeze casting, in which preform design is very important. The preform was manufactured by de-

332 signing the frame with many holes to align the wires parallel. After squeeze casting, the aligned wires can be expected to remain straight due to hydrostatic pressure. AM60 alloy melt was poured into the mold with diameter 80-mm preheated at 300 °C containing the preform. The melt temperature was 750 °C, approximately 100 °C higher than liquidus temperature. The pressure applied for squeeze casting was 100 MPa, the pressure delay time was 7 seconds, and the pressure hold time was 60 seconds. All process was performed under Ar gas and the manufactured composites were cooled in air. The microstructure and main elements of the manufactured composite were analyzed by optical microscopy and scanning electron microscopy (SEM) combined with energy dispersive spectroscopy (EDS). Thin foils for transmission electron microscopy (TEM) were prepared by ion milling under liquid nitrogen cooling. Tensile test at room temperature was performed with an initial strain rate of 8.3 × 10-4/s. The round specimen with gauge length 20 mm and gauge diameter 4 mm was designed. Cross-sectional observations were conducted on the composites to understand the deformation and fracture process. Additionally, Scanning Auger microscopy (SAM) was also used to analyze the elements on the surface of the exposed wires. Sputtering rate for depth profiling was 10Å/min with a Ta2O5 standard.

Results Figure 1(a) shows the optical microstructure (transverse cross-sectional view) of the manufactured composites reinforced with AISI 304 wires. AISI 304 wires are uniformly distributed in AM60 matrix, and there are no defects, which can form by gas trap and/or limited infiltration during the fabrication process. The composites have no reaction layer, but Al-Mn compounds and Mg17Al12 phase, determined by EDS, exist in AM60 matrix as shown in Figure 1(b). There is expected to be a reinforcement element inside the matrix as in the case of AM60 composites reinforced with other metal alloys such as Ni, Al, and Ti [7]. A few nm-sized particles are observed within the AM60 matrix as shown in Figure 2. Selected area diffraction pattern (SADP) analysis shows that the clustered fine particles are pure Fe without reaction with other elements containing AM60 alloy. Such fine Fe particles should be due to Fe diffusion into AM60 matrix during squeeze casting. Mg diffusion into AISI 304 reinforcement could be possible also. Therefore, to check Mg on the surface of reinforcement, composition-depth (sputter time) profiles on the surface areas have been analyzed by using SAM. The composition profiles show that there is Mg up to a few tens nm depth (a few tens min sputter time) from the surface of reinforcement.

a

b

Figure 1. (a) Optical micrograph and (b) SEM micrograph of AISI 304 wire/AM60 composite.

333

a

b

b

Fe particle

100 Figure 2. (a) TEM micrograph of AM60 matrix in the composite and (b) dark field image of fine particles.

Figure 3 shows that the stress-strain curves of AM60 Mg composites are variable with the tensile properties and the volume fraction of AISI 304 wire, respectively. As shown in Figure 3(a), high strength of the composite reinforced with as-drawn wire (19 Vol. %) is due to good microstructure as mentioned above. As the volume fraction of the wire decreases down to 13 Vol. %, the strength decreases but the elongation increases in the composite, respectively. In the case of annealed wire (700 oC for 1 min) reinforced composites as shown in Figure 3(b), their strength decreases but elongation increases in comparison with those of the composite reinforced with as-drawn wire. In the case of annealed wire (700 °C for 3 min) reinforced composites as shown in Figure 3(c), their strength and elongation have the exactly same trends in comparison with those of the composite reinforced with annealed wire (700 °C for 1 min). The result shows that both matrix and reinforcement are uniformly elongated at the same time until debonding at the interface. Tensile properties and fracture type of the composites are strongly dependent upon the elongation of each wire. The composite reinforced with lower elongated wires than matrix has wire fracture, in contrast the composite reinforced with higher elongated wires than matrix has wire debonding and pull-out.

334

(a)

19 vol. %

(b) 13 vol. %

19 vol. % 13 vol. % 9.1 vol. %

(c) 17 vol. % 13 vol. % Figure 3. Stress (MPa)-strain (%) curves of AM60 Mg composites reinforced with (a) as-drawn wires, o (b) wires annealed for 1 min at 700 C, and (c) wires o annealed for 3 min at 700 C.

Discussion The present investigation represents that the stress-strain curve of the composites is variable with those of the reinforcement. To the best of our knowledge, it has been only reported that stress-strain curve of metal matrix composite varies with mechanical properties of matrix under fixing reinforcement. Figure 4(a) clearly shows how stress-strain curve of the composite varies with mechanical properties of reinforcements under fixing all other conditions such as volume fraction, interface condition, matrix, etc. The elongation of the composites is impressive, different from that of the composite reinforced with ceramic fiber, showing the possibility of overcoming the basic problem of metal matrix composites. In conclusion, the best-combined tensile properties of the present composites represent tensile strength of 480 MPa and elongation of 5.0 % in the composite reinforced with the 19 vol. % wires annealed at 700 °C for 1min. Additionally, Figure 4(b) shows the plots obtained from the elongation of the composites as a function of the volume fraction of wire in each composite. Specifically, the composite reinforced with the wires annealed at 700 °C for 1 min has the relatively lower decrease rate of the elongation than the other composites. Furthermore, the elongation of this composite shows much higher value than that of the wire as reinforcement. Such phenomena result from the unique plastic deformation behavior of the composite.

335

(a)

(b)

As-drawn

o

Annealed at 700 C for 3min.

Annealed at 700oC for 1min. Annealed at 700oC for 1min. o

Annealed at 700 C for 3min. As-drawn

Figure 4. (a) Stress (MPa)-strain (%) curves of AM60 Mg composites varying with the wires (13-vol. %) and (b) plots of elongation of composite versus volume fraction of fiber in each composite.

Fracture strain of metal matrix composite reinforced with ceramic fiber is similar to that of the reinforcement, below 1 % [8]. However, the elongation of the present composite is much higher than that of wire due to unique characteristics. Firstly, metallic wires are uniformly elongated inside the matrix. Secondly, fracture type of metallic wire is different from that of ceramic fiber. Metallic wire has a necking (reduction of cross-sectional area) cup & cone [7] like ductile fracture mode, even below 1% elongation. Thirdly, considered as the most important point, the metallic wire has unique plastic deformation inside the metal matrix. From the optical cross-sectional view of tensile specimen, the necking of wire is observed at internal area remote from the final fracture surface as well as the final fracture surface as shown in Figure 5(a) and (b), respectively. More number of necking is observed in the composite with high strength and elongation by using the wire annealed at 700 °C for 1 min as reinforcement as shown in Figure 5(c). Such phenomenon enables the total length of wire to be extended more than the elongation of wire, so as to improve the elongation of metal matrix composite.

a

c

b

200μm

400μm

Figure 5. Cross-sectional optical micrographs of the tensile specimen (longitudinal direction): (a) final fracture surface, (b) internal area remote from final fracture surface of as-drawn wires reinforced composite, and (c) o multiple-necking of AISI 304 wire annealed at 700 C for 1 min in AM60 Mg composite.

336

Summary The microstructural characteristics of AM60 Mg composites reinforced with AISI 304 wires shows that squeeze casting for metal/metal composite is a sound process without casting defect, misaligned wire, and damaged wire. The microstructural analysis of the composite shows no intermediate phase at the interface and fine Fe particles in the matrix. The tensile properties of the composites are dependent upon volume fraction and mechanical properties of wire. The best combination of tensile properties represents tensile strength of 480 MPa and elongation of 5.0 % in the composite reinforced with wires (19-vol. %, annealed at 700 °C for 1min). Multiple-necking of wire can not only extend application of ROM up to plastic deformation region, but also improve the elongation of the composite. Multiple-necking, considered as one of the mismatch accommodation processes between two phases, is both scientific interest (mechanism and modeling) and practical interest (effective condition on manufacture).

References [1] [2] [3] [4] [5] [6] [7] [8]

M. R. Pinnel, A. Lawley, Metall. Trans. A. 1970, 1, 1337-1348. R. B. Bhagat, Metall. Trans. A. 1985, 16, 623-628. S. Y. Kim, C. Y. Hyun, J. Kor. Inst. Met. & Mater. 1987, 25, 351-359. S. M. Lee, H. J. Lee, C. P. Hong, J. Kor. Inst. Met. & Mater. 1990, 28, 225-234. F. Boland, C. Colin, F. Delannay, Metall. Trans. A. 1998, 29, 1727-1739. S. Sun, M. Pugh, Mater. Sci. Eng. A. 2001, 300, 135-141. J. J. Kim, Ph.D. Thesis, POSTECH 2000. C. G. Lee, Sunghak Lee, Dongil Kwon, I. M. Park, Scripta Metall. 1993, 28, 949-954.

337

Thermal Strain in Mg Composites Alexandra Rudajevová, Pavel Lukáč Department of Metal Physics, Charles University, Prague

1

Introduction

Magnesium composites have found many applications in various branches of industry. They exhibit excellent thermal properties such as low coefficient of thermal expansion, relatively high thermal conductivity, and improved mechanical properties as high specific strength. The use of magnesium composites requires a better understanding not only their mechanical properties but also their thermal expansion characteristics. The thermal expansion behavior of MMCs influences the mechanical behavior of the composites (e.g. the stability, geometrical changes and changes in mechanical properties). It is expected that the coefficient of thermal expansion (CTE) of the magnesium composites will be influenced by the kind, volume fraction and form of the reinforcement and by the matrix. The aim of this work is to investigate the thermal expansion behavior of three different magnesium composites. Commercial pure Mg, Mg4Li alloy and magnesium alloy Mg - 6Zn - 3Cu (ZC63) reinforced with short fibers of Al2O3 (Saffil) were used.

2

Experimental

The samples of Mg – 10 vol.% Saffil fibers and Mg4Li - 10 vol.% Saffil fibers composites were prepared by pressure infiltration of Saffil fibers preform in an autoclave under an argon pressure (up to 6 MPa) at temperatures 888- 908 K. These materials were prepared in the Institute of Materials and Machine Mechanics, in Bratislava (Slovakia). Reinforced ZC63 - 25 vol.% Saffil and unreinforced ZC63 alloy were produced by squeeze casting in the Technical University of Clausthal (Germany). The linear thermal expansion of the as-cast samples was measured in argon atmosphere using the Netzsch 402E dilatometer, over a temperature range from room temperature to 380 °C at heating and cooling rates of 5 K/min. The volume concentration of the Saffil fibers was determined from the density measurements.

338

3

Results 5

1.run 2.run mix. rule pure Mg Saffil

0.010

0 -5 Δd/lo [10-4]

relative elongation

0.008

0.006

0.004

-15

0.002

0.000

-10

1.run 2.run Mg-1.run Mg-2.run

-20 -25 0

100

200

300

400

o temperature [ C]

Figure 1. Temperature dependence of the relative elongation for the Mg-10vol.% Saffil composite

0

100

200

300

400

temperature [oC]

Figure 2. Temperature dependence of the thermal strain for Mg and Mg4Li composites

The temperature dependence of the relative elongation of pure Mg and composite Mg with 10 vol% Saffil (two runs), are shown in Fig. 1. The dotted curve in Fig. 1 was obtained from the relative elongation of Mg and Saffil using the rule of mixtures. The relative elongation of Saffil was obtained from [ 1]. Fig. 2 shows the temperature dependence of the thermal strain Δd/lo for Mg and Mg4Li composites. This parameter is defined by the following relation: ⎛ Δl ⎞ Δd ⎛ Δl ⎞ =⎜ ⎟ −⎜ ⎟ lo ⎝ lo ⎠ meas. ⎝ lo ⎠ mix.rule

(1)

where (Δl/lo)meas. is the measured relative elongation of the composite and (Δl/lo)mix. rule is the relative elongation of the composite calculated by the rule of mixtures. The thermal strain Δd/lo shows the departure of the composite behavior from the case where no thermal stresses are present. Fig. 2 reveals that after the first thermal cycle the irreverisible changes of the sample length occur. While there is irreversible contraction for Mg composite (10 μm), m), the extension was found in Mg4Li composite (19μm). m). In Fig.3 the temperature dependence of the thermal strain is shown for the composite ZC63-25 vol.% Saffil for the maximum temperatures of the thermal cycle 380 °C (irreversible contraction after the first run was 8 μm) and 530 °C (irreversible contraction 95 μm). A sudden increase of the Δd/lo above 450 °C occurs as result of the phase transition in alloy.

339

4

Discussion

During cooling after fabrication of MMCs, a large difference in the coefficients of thermal expansion between the metallic matrix and the ceramic reinforcements causes large residual thermal stresses. These thermal stresses have the highest values at room temperature. During following heating, thermal stresses are released by elastic and plastic deformation. The development of the thermal strain for Mg and Mg4Li composites for two thermal cycles is shown in Fig. 2. It can be seen that that dilatational characteristics of Mg and Mg4Li are different. The two main differences between both composites are important:

-4 Δd/lo [10 ]

-4 Δd / lo [10 ]

1) While the irreversible contraction occurs in Mg composite after the first run, the irreversible extending occurs in Mg4Li composite. 2) While hysteresis of the temperature dependence of the Δd/lo is localised above 250 °C for Mg4Li composite, for Mg composite it was found in the whole temperature range studied. The thermal strain for the ZC63 composite has the similar character as Mg composite (Figs. 2 and 3). The different behavior of Mg4Li composite in comparison to pure Mg and ZC63 composites is very probably caused by the strength interfacial bonding. During infiltration of the Saffil preform by the Mg4Li alloy at high temperature, lithium may intereact with the Saffil fibres. The interaction of Li and Mg with the δ-Al2O3 (Saffil) was analysed 0 in details by Kúdela [2]. The better adhesion MgLi alloys to Al2O3 fibers in comparison with the system Mg-Al2O3 is due to the -20 different ability of the ions Li+ and Mg+ to enter into the vacant cation interstitial positions of δ-Al2O3 lattice. The topotactic -40 entrance Li+ into δ-Al2O3 lattice turns the 1.run oxid-oxid structure in the interface and -60 therefore, binding between the fiber and the 0 matrix is stronger than in Mg composites where entrance Mg+ into δ-Al2O3 lattice is -20 not possible. The presented results show that Li considerably influences the chemical and phase composition of the interface and -40 therefore the thermal stress localised at the 2.run interface. From our results we may assume -60 0 100 200 300 400 500 600 that Zn and Cu (in ZC63 composite) does not influence conditions in the interface in temperature [oC] comparison with the Mg composite. The results for Mg and ZC63 composites Figure 3. Temperature dependence of the thermal may be explained by the model presented strain for ZC63 composite in a previous paper [3]. According this model the shear stress due to the difference in the CTE between both components has a possitive starting value at room temperature. Heating the composite slowly, the shear stress decreases firstly elastically, remaining below the relevant plasticity limit (Fig. 4). At a certain temperaure when the shear stress reaches the plasticity limit, creep starts near the

340 interface. Switching the heating to the cooling during creep regime, the representative shear stress is firstly raising elastically. If the limit of reverse creep is reached, the shear stress increases with decreasing temperature up to the temperature where the complete elasticity occurs. We suppose that the temperature dependence of the thermal strain in the second thermal cycle indicated relaxation of the thermal stress by only elastic deformation in the whole temperature range. When the thermal stress in the matrix reaches the yield stress, plastic deformation occurs near the interface. The thermal strain decreases with increasing temperature more rapidly than in case when the stress in the matrix would be lower than yield stress. It is supposed that plastic deformation induces the sources of the internal stresses which modified the elastic stress field in the matrix. The experimental consequnce of this process is a change in the temperature dependence of the dilatational characteristics of the composites Mg4Li alloy has nearly the same Emodulus and limit of plasticity as pure Mg Figure 4. Variations of the maximum shear stress [4]. The dilatational characteristics are howduring heating and cooling (skech) ever different. The proposed model or various relations used for the calculation of the thermal stress (e.g. Careno-Morelli relation et al. [5]) cannot elucidate the different behaviour of both composites because they does not consider the changes of the physical properties of matrix near the interface when chemical reactions between matrix and fiber occur.

5

Conclusions

The thermal stesses in all Mg composites investigated influence the thermal expansion characteristics. The relative elongation in the first thermal cycle is different from that in the second one. While in the composites based on pure Mg and ZC63 alloy the permanent reduction was found, in Mg4Li composite the permanent extension of the sample was found after the first thermal cycle. While the hysteresis in Mg and ZC63 composites was found in the whole studied temperature range, the hysteresis in Mg4Li composite was found above 250 °C. This different behavior of the studied composites is ascribed to the interaction of Li with fibers.

341

Acknowledgements The authors acknowledge financial support from the Grant Agency of the Academy of Sciences under Grant A2112302 and from the Grant Agency of the Czech Republic under Grant 106/03/0901.

References [1] S. Chang, S. Lin, M. C. Flemings, Metall. Mater. Trans. 2000, 31A, 291. [2] S. Kúdela, Fibre-Matrix Interaction in Mg Composites, PhD Thesis, Institute of Materials and Machine Mechanics, Slovak Academy of Sciences, Bratislava, 2002. [3] A. Rudajevová, J. Balík, P. Lukáč, Sci. Eng. Comp. Mater. 2000, 9, 11. [4] I. I. Gurjev, F. M. Elkin, V.G. Kudrjašov in: Struktura i svojstva legkich splavov, “Nauka“, Moscow, 1971. [5] E. Careno-Morelli, S. E. Urreta, R. Schaler, Acta. Mater. 2000, 48, 4725.

342

Synthesis of TiC Dispersed Composite in Mg-alloy Matrix D.W. Lee1, S.V. Aleksandrovkii2, V.M. Sizyakov2, D.V. Kutsenko2 1 2

Korea Institute of Machinery and Materials. Changwonn, Kyungnam, Republika Korea. St. Petersburg Mining Institute (Technical University), Russia.

Abstract Carbide dispersed metal matrix composite (MMC) was produced successfully by the feeding the mixed solution consisted titan and carbon chlorides into liquid magnesium alloy. It was revealed that the produced MMC had the magnesium alloy matrix with the compositions of 7,5% Al, 0,2 ∼ 0,8% Zn, 0,15 ∼ 0,5 Mn, in which the particles of titanium carbide were uniformly reinforced with a needle shapes. A armouring phase of particles consisted of titanium and aluminum. X-ray diffraction pattern has shown the clean peaks of titanium carbide. High mechanical strength and young,s modulus were obtained in the produced composite materials.

1

Introduction

Problems of new composite materials production have been pointed at the end of this century, in fields of nuclear energetic, air-space industry, machinery needs these materials. Steel is no more popular now as last time. So composite materials have been regulated to be developed [1]. The composite materials, based on the light metal is more interesting materials recently. It is known that MMC has shown good properties of the most part as the famous starting metals and compounds. Aluminum-magnesium alloys have used as matrix metal for MMC. And different type of materials: carbides, oxides, nitrides and etc. have used as hardening phase. If there are no suitable metals, which should be used, as the serious function, MMC can contribute for good application.

2

Theory grounding

There are two types of MMC. Here type is depend on the production process of the MMC [2]: – “Natural” MC is formed as result of the chemical reaction process, new phase precipitation and starting reagent reduction process. Matrix surface and hardening phase do

343 not contain impurities and have good properties of hardness and elastic modulus. Phases of “natural” MMC are in thermal dynamical equilibrium. – “Synthetic” MC are produced by two phase mechanical mixing, with the following consolidation, heating process and extrusion. Mixed phases usually have no thermodynamical equilibrium. Chemical reactions are possible during the production process because of that. Synthetic MMC are various, but particle surface can content the impurity such as oxide. So synthetic MMC are not so strength as “natural” because of these contaminations. MMC process has some difficalty due to the complicate production process, and the cost is higher than that one of non-reinforced materials. Main technical difficulty is thermal and chemical compatibility of matrix and dispersoids. In particular, Al cannot wet titanium carbide well. At direct mixing of the starting reagents Al and carbide it is difficult to recover the good results during the production process. Production technology of metal matrix composite (MMC) by the metal thermal reduction method has been positively investigated. The process can be characterized with the following reaction and operations: TiCl4 + C2Cl4 + Al(Mg) → TiC – Al(Mg) Al(Mg), Mg(Al) Ar TiCl4 CCl4

reduction 800 – 900 °C

N2

mix 800 – 600 °C

TiC – Al(Mg), TiC – Mg(Al) TiCN – Al(Mg)

The physico – chemical properties and application of MMC production have been researched to find the best composite to optimize the production process. The thermodynamical analysis has shown the good possibility for production of titanium carbide, titanium nitride and carbide nitride forming. Kinetic characteristics of the heterogeneity process in the system: refractory compounds – chlorides of refractory and alkaline metals can be detected by thermal stability of starting and final compounds. Magnesium thermal reduction process with using the mixtre of o carbon tetrachloride and titanium chloride is performed at 830-1000 C. This reduction process has a diffusion mode [3-5]. Kinetic mode is realized at the interaction process of

344 magnesium and chloride, temperature is higher that 1000 °C. It can be explained by the full decomposition of carbon tetrachloride and formation of the dispersed carbon out side of the reaction zone. It has provided the synthesis of non - stoicheometric titanium carbide (approximate composition is TiC0, 6 - TiC0, 9).

3

Technique of process

The mixture of titanium chloride and carbon chlorides was poured on the surface of mixed magnesium and aluminum melt to obtain the MMC. The temperature of this process was hold at 800-900 °C. After the end of this process the oproduced product were cooled down. Final temperature of the mixing process has a strong effect on the homogeneity of the obtained product. For example, optimal temperature of magnum alloy ML-5 is 590 oC. The microstructure of obtained MMC was investigated. Composite material based on the alloy ML-5(Al 7,5-9, Zn 0,2-0,8 and Mn 0,15-0,5%) were represent as magnesium matrix, reinforced by the range of titanium carbide grains(fig.1).

Figure 1. Microstructure of KMM based on Mg alloy ML5 and mapping morphologies of Al(2) and Ti(3)

Local part of this sample has the high amount of titanium carbide. It can be explained by the high temperature effect of the final mixing process - 650 °C. Mapping analysis has shown the presence of Al and Ti in reinforce particle. The presence of the Mn wasn't shown. Line scanning confirms the presence Al and Ti reinforce particle (fig.2). The EDS in SEM microstructure has also confirmed the present of titanium and aluminum inside

345 the reinforced particles; synchronously, the other elements of Mn and Fe were moved to the particle surface from matrix. X-ray analysis shows the titanium carbide phase exactly. MMC based on aluminum alloy AlMg61 (Mg 5,5 6,5, Mn 0,7-1,1, Zr 0,05-0,15) have shown the presents of reinforced particles with the different configuration (fig.3). Mapping test has shown the present of carbonized titanium particle that has the round and not correct forms. Fe and Mg have the particles of acute-angled structure. Aluminum base MMC were scanned (fig.2).

Figure 2. Line profile on the cross sectional particle for Mg base MMC (A) and Al base MMC (B). x-axis unit [mkm], I-intensively (impulse).

Figure 3. Microstructure of Mc based on Al alloy AMg61(1) and mapping morphologies of Mn(2) and Ti(3).

346 This scanning has confirmed the titanium and Mn presents in reinforced particles in particles. It is very important to mention that we can observe the aluminum content. This content is decreasing by the particle surface. X-ray analysis of the obtained samples has shown the titanium carbide present. Phase of titanium carbide and Ti3AlC is shown very weakly. In reinforced TiC particles, fixed carbon content changed from 15 to 19%.

4

Conclusion

Executed research work allow to present the following stage of MMC production process: – TiCl4 and CCl4 reduction with obtaining of the Ti and C elementary particles and titanium carbide forming with near stoicheometric composition of TiC; close contact of the first stage formation components (Ti and C) promotes the activity of this process. Distance between the particles can be related with the solution homogeneity of starting titanium chloride and carbon particle molecules, which are unlimited solute in each other. Matrix refining (Al-Mg alloy) and magnesium chloride formation is going on at the same time. – MMC production; when TiC can be produced identically during in-suty process the surface of reinforced phase and matrix are not impureted and have a high activity. MMC has reinforced particles of endogenous origin. It is possible to propose that TinAl and triple compound of TimAlC provide the adhesion force increasing on the border between the matrix and reinforced particles in MMC production base on aluminum alloys. It can be confirmed by aluminum present in reinforced particles (fig. 2). In consequence of interaction with described process and following heterogeneous system TiCl4CCl4-Mg-Al-MgCl2-C-Ti-TiC produced MMC can have high hardness and elastic modulus. Phisico-chemical and technological foundation for the production process of new highstable construction materials, which based on magnesium thermal reduction process of rare metal chloride compound, have been developed. Industrial equipment for magnesium thermal reduction process can be used.

5

Reference

[1] D.R. Hiatt. JOM. 1990, 42, № 3, 32-35. [2] E.Z. Fine Morris. Metallk. 1993,84, № 4, 282-285. [3] S.V. Aleksandrovkii, V.M. Sizakov, A.H. Ratner, D.W. Lee. J. Materials Processing and Manufacturing science. 2001, 9, 303-309.

347 [4] S.V. Aleksandrovkii, V.M. Sizakov, A.H. Ratner. Titanium 99. Science a Technology. Prodeegings of the 9th World Conference on Titanium. 7-11 June 1999, StPetersburg, 1834-1839. [5] S.V. Aleksandrovkii, D.W. Lee, V.G. Gopienko. New methods of production refractory titanium compounds. Ruda and Metalls. St-Petersburg, 128.

Influence of Fibre Preform Permeability on Infiltration of Magnesium-Zinc Base Alloys M.S. Yong*, R.I. Temple† and A.J. Clegg** * Singapore Institute of Manufacturing Technology, Singapore **Wolfson School of Mechanical and Manufacturing Engineering, Loughborough University, UK

1

Introduction

A number of researchers have evaluated the feasibility of producing fibre reinforced magnesium alloys [1-4]. The results showed that the mechanical properties of magnesium alloys can be significantly enhanced with fibre reinforcement. Ha [3] and Chadwick [4] investigated the effect that various reinforcement elements had on the mechanical properties of magnesium alloy AZ91. They reported that mechanical properties such as hardness, fatigue strength and creep resistance at elevated temperature were greatly improved for the reinforced alloy. Typically the fatigue strength of AZ91, reinforced with 16% vol. Saffil fibre, at 180 °C was increased by 100%. Despite the significant advantage of using magnesium MMC for lightweight and high strength applications, very little work has been conducted to investigate the relationship between permeability and the infiltration of fibre preforms. Prior to such an investigation, appropriate fibres and binders must be selected. This is because the choice of fibres and binders is known to have significant influence on the MMC. In particular, the main criterion that influences the selection of fibre type is compatibility with the matrix. Two fibres that are compatible with magnesium are Saffil and carbon [5]. Silica and alumina-based binders are widely used in preform production, mainly due to their high temperature properties [6]. However, there are concerns about chemical reactions between the magnesium alloys and the silica (SiO2) coating of the reinforcement [7]. Hence the influence of these fibres and binders, when used with magnesium alloy, was evaluated by examining the mechanical properties and microstructures of the cast samples.

2

Experimental Procedure

The magnesium alloy used in this investigation was Mg-4.2%Zn-1%RE, hereafter referred to as RZ5DF alloy. The composites were produced using a squeeze infiltration technique. Melting was conducted using the fluxless method with argon gas as a protective cover. The die was coated with boron nitride suspended in water to protect it from excessive wear. An applied pressure of 80 MPa and a preform temperature of 600 °C were used to squeeze infiltrate the alloy. A preform temperature of 600 °C or above was necessary to minimise resistance to magnesium infiltration at the preform surface [8]. Five preform systems were selected to investigate their suitability when squeeze cast with RZ5DF alloy. They were (i) 20% vol. fraction Saffil fibres with 5% silica binder, (ii)

349 20% vol. fraction Saffil fibres with 5% alumina binder, (iii) 14% vol. fraction Saffil fibres with 5% silica binder, (iv) 12% Saffil + 9% carbon fibres with 5% alumino-silicate binder, and (v) 20% vol. fraction carbon fibres with 5% alumino-silicate binder. The study aimed to establish the extent of improvement that each preform system provided to the magnesium alloy. Comparisons were made between the mechanical properties and microstructures of the MMCs fabricated from the five preform systems. Standard procedures were used for the preparation of specimens and the evaluation of mechanical properties [8]. The permeability meter, test set-up and the preforms used for the tests are shown in Figures 1 and 2.

Figure 2. Photograph of different preform systems shaped to evaluate their permeability. The preforms shown (from left to right) are: (i) (ii) (iii) (iv) Figure 1. Permeability meter and test components. The assembly of test components was fitted on the orifice (top portion) of the meter.

20% vol. fraction Saffil fibres with 5% silica binder 20% vol. fraction Saffil fibres with 5% alumina binder 14% vol. fraction Saffil fibres with 5% silica binder 12%Saffil + 9%carbon fibres with 5% alumino-silicate binder (v) 20% vol. fraction carbon fibres with 5% alumino-silicate binder

3

Results and Discussion

3.1

Permeability Measurement of Fibre Preforms

The quality of preforms supplied by the manufacturer was examined by SEM. This ensured sound preforms (i.e. no fibre fracture) prior to squeeze infiltration with the RZ5DF alloy. Such pre-examination was necessary, as fractured fibres are detrimental to the overall MMC mechanical properties. The examination revealed minimal to zero fibre fracture. This eliminated poor quality preforms as a potential cause of poor composite mechanical properties. A study of fibre preform permeability was conducted to evaluate the effect of permeability on the degree of infiltration. The property of permeability considered in this research refers to the ease with which gases will permeate through a porous fibre network. This property is important as it influences the efficiency with which the melt infiltrates the preform. During infiltration, the metal must first displace the air from within the preform before it can replicate the cavity shape. Failure to do so will result in porosity in the MMC casting. Table 1 summarises the results from the study.

350 The permeability results show that the carbon fibre preform was the most permeable, followed by the preforms containing 14% vol. Saffil fibres, 20% vol. Saffil fibres with silica binder and 20% vol. Saffil fibres with alumina binder. The least permeable was the preform containing 12% vol. Saffil + 9% vol. carbon fibres. Table 1. Results of Permeability measurements for different preform systems. (The greater the permeability number the more permeable the preform is). Measurements were taken over an area of 18mm X 40mm X 15mm thickness, using a large orifice.

No. 1 2 3 4 5

Preform Fibres 20% Saffil 20% Saffil 14% Saffil 12% Saffil + 9% Carbon 20% Carbon

Binders Silica Alumina Silica Alumino-silicate Alumino-silicate

Permeability Number Average of three Readings 180 160 247 83 350

The higher permeability of the carbon fibre preforms may be attributable to their larger fibre diameter (8μm) when compared to Saffil fibre (3μm). This is similar to the permeability of bonded sand moulds, large sand grains have large pores between them thus allowing air to permeate more readily. Conversely, small grains (consider smaller fibre diameters in this investigation) have smaller pores or voids which creates a more difficult and tortuous route through which the air must pass. Larger diameter fibres provide high permeability but a weaker MMC. There are two reasons for this. Firstly, larger diameter fibres have a greater possibility of containing crystalline imperfections thus producing a weaker fibre element for the MMC. Secondly, for a given fibre volume, a smaller diameter fibre has more surface area in contact with the matrix than a large diameter fibre. Based on the mechanics of reinforcement, the MMC strength is dependent on the fibre diameter and its surface area. The use of smaller diameter fibres has, therefore, contributed to the improvement of the MMC strength. Hybrid preforms that contain two or more sizes of fibres, will have a different permeability value. A mixture of fibre sizes and distribution tends to result in the filling of voids between larger fibres (e.g. carbon) by the smaller fibres (e.g. Saffil), thus reducing the overall permeability. The low permeability number obtained in the study for the 12% Saffil + 9% carbon volume fraction fibre preform supports this assumption. 3.2

Tensile and Hardness Properties

The tensile properties of the five composites were compared. The results showed the 14% vol. fraction Saffil fibre preform to be the most promising, in terms of the ease of production and maximum value to strength performance. Figure 3 summarises the tensile properties of the five composites. The highest UTS values (259 to 284 MPa) were obtained from Saffil reinforced composites and the lowest for the composites containing carbon fibres (181 to 233 MPa). The lower UTS value for the carbon reinforced RZ5DF alloy was attributed to the weak bond between the carbon fibres and the magnesium matrix. The wetting characteristics of carbon fibres play a major role in the strength of the composite. Carbon fibres have poor wetting characteristics with magnesium that results in an increased resistance to infiltration. As a consequence the high applied pressure of squeeze

351 casting caused fracture of the carbon fibres. This resulted in ineffective fibre reinforcement as the fibre lengths were below the critical value resulting in lower tensile properties. The hardness values of the magnesium matrix composites produced using different preform systems were compared. The results show a complex correlation between several factors that influence the hardness value of the magnesium matrix composite. Figure 4 shows the material hardness along the ‘longitudinal’ and ‘transverse’ directions for the five preform systems. A comparison of hardness values shows the large differences between the systems. For instance, the RZ5DF alloy reinforced with 20% vol. Saffil + alumina binder provided an average hardness value of 80 HRB, compared to 62 HRB for the RZ5DF alloy reinforced with 20% vol. carbon + alumino-silicate binder. The factors that may have caused variations in hardness include (i) incorrect selection of process parameters and/or preform system causing deformation of the fibre preform during infiltration; (ii) higher fibre content increases the hardness of the RZ5DF alloy because more stiff and hard fibres are present in the composite; and (iii) fibre segregation. A detailed explanation of the causes of hardness variation can be found in the author’s thesis [8]. 3

150

284

266

259 233

100

181

50

2

% Elongation

200

1.5 1

% Area Reduction

2.5

250

Percentage (%)

Ultimate Tensile Strength (MPa)

300

0.5

0

0

20% Saffil + 20% Saffil + 14% Saffil + 12% Saffil + 20% Carbon silica binder alumina silica binder 9% carbon fibres binder fibres

20% Saffil + 20% Saffil + 14% Saffil + 12% Saffil + 20% Carbon silica binder alumina silica binder 9% carbon fibres binder fibres

85

90

80

85

Rockwell Hardness (HRB)

Rockwell Hardness (HRB)

Figure 3. Tensile properties for the five preform systems after infiltration with RZ5DF alloy. Cast with constant pouring temperature of 750 °C and die temperature of 250 °C.

75 70 65 60 55 50 45 40 1

2

3

4

5

6

7

8

9

10 11 12 13 14 15

Longitudinal dire ction along the Spe cime ns

14% Saffil

80 75

20% Saffil w Silica B.

70 65

20% Carbon

60 12%Saffil 9%Carbon

55 50

20% Saffil w Alumina B.

45 40 1

2

3

4

5

6

7

8

9

Trans vers e direction along the Specimens

Figure 4. Material hardness along the ‘longitudinal’ and ‘transverse’ direction for the five preform systems. Cast with constant pouring temperature of 750 °C and die temperature of 250 °C.

3.3

Metallography

Microstructural examinations of magnesium composites reinforced with (i) 20% vol. Saffil + silica binder, (ii) 20% vol. Saffil + alumina binder and (iii) 14% vol. Saffil + silica binder showed sound cast structures. However, there was evidence of fibre cluster-

352 ing with the higher volume fraction of fibres. Microstructural examinations of magnesium composites reinforced with 12% vol. Saffil + 9% vol. carbon revealed significant porosity in the region near the edge of the fabricated composite. Typical examples of this porosity are shown in Figure 6. Microstructural examination of the carbon reinforced magnesium alloy composite showed significant fibre fracture throughout the composite structure, seen in Figure 5. This indicates ineffective fibre reinforcement, as the fibre lengths have fallen below the critical value. The importance of critical fibre length for fibre reinforcement has been emphasised by Anderson et. al. [10] and Harris [11]. As a consequence of shorter fibre length, lower tensile properties were obtained for the cast carbon reinforced composite. The microstructural examinations of the five MMC systems revealed that Saffil preform systems showed sound cast structures, i.e. no porosity. However, evidence of fibre clustering was seen with a higher (20%) volume fraction of fibres. This may be attributed to the fact that denser preforms present greater resistance to full infiltration by the melt. The hybrid composite containing 12% vol. Saffil + 9% vol. carbon exhibited significant porosity within the composite, which can be attributed to the low permeability of the hybrid perform. The magnesium-carbon reinforced composite exhibited significant fibre fracture throughout the composite structure, which led to ineffective fibre reinforcement and the lower tensile strength of 181 MPa. This is in comparison to the higher UTS value (266 MPa) of the RZ5DF alloy with 20% vol. Saffil + silica binder, a difference of 85 MPa.

Figure 5. Micrograph of the RZ5DF-20% vol. carbon reinforced composite, showing the common feature of fibre fracture.

4

Figure 6. Micrograph of the RZ5DF-12% vol. Saffil + 9% vol. carbon reinforced composite, showing the presence of shrinkage porosity in a particular region of the composite specimen.

Conclusions

1. A fibre preform permeability test was used to assess the suitability of preforms for infiltration. However, although there was a correlation between permeability and infiltration capability, this was not the only factor of importance. Of equal importance were the wetting characteristics of the preform and matrix combination. 2. The fibre preform permeability test was justified as one mechanism to identify the ease of infiltrating the metal into the preform. Quality assessment of the preform prior to casting is important because it is necessary to ensure preform permeability consistency from commercial production. Hence, the permeability test may be used as a simple quality control test in industry for the evaluation of preform consistency provided by the supplier(s).

353 3. Fibre preforms containing carbon fibres were only partially infiltrated by the RZ5DF alloy. This incomplete infiltration was attributed to poor wetting between the carbon fibres and magnesium, since both elements are known to be thermodynamically inert to each other. The ineffective fibre reinforcement resulted in lower tensile properties. 4. The preform system consisting of 14% volume fraction Saffil with silica binder was found to have the best value to strength ratio. The low volume fraction of fibres allowed greater ease of infiltration and the good wetting between the magnesium and the Saffil fibres reduced the risk of fibre clustering during squeeze infiltration.

References [1] Purazrang, K., Abachi, P. and Kainer, K.U., “Investigation of the Mechanical Behaviour of Magnesium Composites”, Composites, 1994, Vol. 25, No. 4, 296-302. [2] Kamado, S. and Kojima, Y., “Microstructures and Tensile Properties of Mg-Zn Based Alloy Composites Reinforced with δ-Al2O3 Short Fibre and 9Al2O3.2B2O3 Whisker”, Science and Engineering of Composite Materials, 1997, Vol. 6, No. 3, 159-167. [3] Ha, T.U., “Squeeze Casting of Magnesium-Based Alloys and Their Metal Matrix Composites”, PhD. Thesis, University of Southampton, 1988. [4] Chadwick, G.A., “Squeeze Casting of Magnesium Alloys and Magnesium Based Metal Matrix Composites”, Magnesium Technology Conf. Inst. Metals, 3-4 Nov 1986, 75-82. [5] Rohatgi, P., “Advances in Cast MMCs”, Advanced Materials & Processes, Feb 1990, 39-44. [6] Clyne, T.W. and Withers, P.J., “An Introduction to Metal Matrix Composites”, Cambridge University Press, 1993. [7] Inem, B. and Pollard, G., “Interface Structure and Fractography of a MagnesiumAlloy, Metal-Matrix Composite Reinforced with SiC Particles”, Journal of Material Science, 1993, Vol. 28, 4427-4434. [8] Yong, M.S., Process Optimisation of Squeeze Cast Magnesium-Zinc-Rare Earth Alloys and Short Fibre Composites’, PhD Thesis, Loughborough University, 1999. [9] Ottinger, O., Grau, G., Winter, R. and Singer, R.F., “The Effect of Aluminium Additions on the Interfacial Microstructure and Mechanical Properties of C/MgComposites”, The Tenth International Conference on Composite Materials (ICCM10), 1995, Vol. VI, 447-454. [10] Anderson, J.C., Leaver, K.D., Rawlings, R.D. and Alexander, J.M., “Materials Science”, Chapman & Hall, 4th Edition, 1990, 326-353. [11] Harris, B., “Engineering Composite Materials”, Institute of Metals, London, 1986.

354

Magnesium Integral Foams – Production and Properties Markus Hirschmann, Matthias Lamm, Carolin Körner, Robert F. Singer Department of Materials Science, Lehrstuhl Werkstoffkunde und Technologie der Metalle (WTM), Friedrich-Alexander-Universität Erlangen-Nürnberg, Erlangen, Germany

Introduction Metal foams offer a high potential for light-weight construction. Their outstanding properties are the low density in combination with extraordinary mechanical, electrical, thermal and acoustic properties. Especially the high bending stiffness at low specific weight offers a great potential for light-weight construction. Meantime a wide range of production routes exists especially for aluminum foams [1]. In the case of magnesium only a few foam types are known, like a syntactic foam with embedded ceramic hollow spheres [2], so called GASAR or lotus-type structures made by a solid-gas eutectic solidification process [3], open cellular structures which use polymeric preforms [4] or foams made by infiltration of porous salt preforms [5]. The disadvantages of these processes are often the difficult process control and the high costs which does not allow a cost effective production. Although thermoplastic foams were commercialized in the late 1960s and have meanwhile often replaced the conventional solid parts [6], the production of metal foams, especially magnesium integral foams, in a similar way to plastics by mixing a blowing agent into the melt has not succeeded so far. A reason for this is the difficulty of foaming a metal in opposition to plastics. Liquid metal has a low viscosity, a high surface energy and especially magnesium has a very low thermal capacity which causes a fast solidification. However, the conditions for good foaming are high viscosity and low surface energy. For the first time we succeeded in manufacturing a magnesium integral foam (Fig. 1). This integral or structural foam consists of a dense skin and a porous core with a gradual transition from skin to core which implicates good mechanical properties. We assume that the stabilization of the cellular structure during the solidification process is done by the primary solidified magnesium phases [7]. The parts are produced by a modified casting technique which is suitable for mass production in a cost effective way. In this paper, the emphasis is on the mechanical properties of the magnesium integral foams and the correlation between the density and the mechanical properties.

Figure 1. Cross section of a magnesium integral foam plate (AM60). Thickness 6 mm.

355

2

Experimental

2.1

Foam production

The foam parts were produced with a high pressure die casting machine and an injection molding machine. For this purpose liquid magnesium alloy (AZ91, AM60) was injected into a permanent steel mould with high velocity. To get a porous structure a chemical blowing agent in powder form (MgH2) had to be mixed into the melt (content < 2 wt. %). Therefore, a so called mixing cavity (2) was integrated between the melt pool (1) and the part forming cavity (3) (Fig. 2). In the mixing cavity the blowing agent is swirled into the melt in a turbulent way. As soon as this mixture enters the part forming cavity the blowing Figure 2. Schematic illustration of the agent decomposes because of the elevated temperature casting process: of the melt into magnesium and a large volume of 1: melt pool, 2: mixing cavity, 3: part hydrogen gas. The gas forms a porous structure in the forming cavity kernel of the cast part whereas at the mold walls a dense skin is formed because of the fast solidification. In order to prevent destroying the foam structure a dwell pressure was not applied. Through variation the process parameters integral foams with different density profiles and overall densities can be produced. 2.2

Foam characterization

The micro computer tomography system μ-CT40 from Scanco Medical was used to investigate the inner structure of the integral foams. The foams (sample dimensions: 6 x 5 x 5 mm³) were scanned using a X-ray tube at 60 kV and 133 μA (voxel size: 10 μm). The micro computer tomography is a very practical non destructive technique to measure density, pore size distribution and density profile of light metal foams [8]. Additionally, two-dimensional cross-sections are made which can be reconstructed to threedimensional objects (Fig. 3). 2-dimensional cross-sections

3-dimensional object

relative density = 100 %

6 mm

1 mm

relative density = 70 % 1 mm

relative density = 30 % 1 mm

Figure 3. Two-dimensional crosssections (dark = pores) of an AZ91 integral foam at different relative densities (skin 100%, 70%, core 30%) measured by micro computed tomography (left). The cross-sections can be reconstructed to a threedimensional object (right).

356 2.3

Mechanical testing

Compression tests were carried out in a tensile testing machine to characterize the mechanical behavior of the foam parts at room temperature (sample dimension: 6 x 6 x 6 mm³, test velocity: 0.4 mm/min, test direction: perpendicular to dense skin). The bending stiffness of the integral foams was investigated by three-point bending tests (sample dimension: 55 x 6 x 8 mm³, support width: 40 mm, test velocity: 0.2 mm/min). Additionally, the elastic modulus and the damping behavior of an AZ91 bending sample were determined perpendicular to the dense skin with a resonant frequency and damping analyser system.

3

Results and Discussion

3.1

Density, density profile and pore size distribution

The magnesium integral foams show a dense skin with a thickness of 0.5 – 1 mm (Fig. 4 left). In this area the porosity is nearly zero. A characteristic property is the variation of the density across the part cross section with the lowest density in the core and a rapid density step between the skin and the core. The overall density of the integral foam parts can be adjusted trough variation of the process parameters. A lower overall density results in a lower density in the core. In the majority of cases the density profile is symmetrical, therefore a mechanical anisotropy perpendicular to the cross section can be excluded. With increasing distance from the dense skin, the pores get larger. With decreasing overall density and consequently increasing porosity in the core, the pore size increases, although the majority of the pores stay in a very small size between 100 μm and 500 μm (Fig. 4 right).

Figure 4. Left: Density profile for different overall densities. Right: Pore size distribution for different overall densities.

By varying the ratio of blowing agent and magnesium melt the overall density and the density profile of the foams can be changed. With an increasing amount of blowing agent and a decreasing amount of liquid metal the overall density decreases. 3.2

Compression behavior

The compression tests demonstrate how the stress-strain-curves vary with decreasing overall density (Fig. 5 left). The ultimate strength decreases while the elongation at rupture increases. For densities lower than about 1.1 g/cm³ the typical foam compression

357 behavior appears. Such a curve has different stages (Fig. 5 right). For small compressions the material exhibits an elastic behaviour, this is shown by a short linear relationship of stress and compression strain (1). The elastic region is followed by an extended plateau region which is characteristic for metal foams (2). The plateau results from successive collapse of the foam core. This can be visualized with the help of micro computed tomography (Fig. 6). Having compressed most of the pores the compression of a “compact” material starts and the slope of the curve rapidly increases (3). With decreasing overall density the plateau region broadens and the stresses become lower. A great advantage of this behavior is that a foam can absorb much more energy than a compact material at a given maximum stress level.

Figure 5. Left: Compression curves for different overall densities. Right: Typical foam compression curve with different compression stages (1, 2, 3). The area under the curve is equal to the absorbed energy.

Figure 6. Compression states of a magnesium integral foam (density = 1.21 g/cm³) measured with micro computed tomography (dark = pores). Left: foam before compression, center: after 15% compression, right: after 30% compression.

3.3

Bending behaviour

A possible application for integral foams may be ultra-thin sandwich panels or beams with excellent bending stiffness at low specific weight. If we compare the bending stiffness of a compact magnesium beam with beams made from magnesium integral foam with the same mass, the foam has a higher bending stiffness (Fig. 7 left). With decreasing overall density the relative increase of bending stiffness is linear. At the moment we are able to increase the weight-specific bending stiffness up to 60% for 6 mm panels. In comparison with aluminum or steel the magnesium integral foam beams offer a very promising lightweight construction potential (Table 1).

358

Figure 7. Left: Relative increase of the bending stiffness as a function of the overall density. Right: Weight saving potential of AZ91 foam (density = 1.25 g/cm³).

In comparison with a St12 deep-drawing sheet steel the AZ91 foam (density = 1,25 g/cm³) has nearly twice the bending stiffness of St12 at the same mass. Compared to an aluminum beam the bending stiffness is 65% better, to an AZ91 beam it is 45% better (Table 1). Otherwise comparing the mass reduction at the same stiffness, the AZ91 foam has a 63% lower weight than steel (Fig. 7 right). The weight reduction of a compact AZ91 plate is 10% worse that of the foam, of an aluminum plate it is 15% worse that of the foam. Table 1. Light-weight construction potential of magnesium foam (mass = const.)

Beam thickness Improvement of bending stiffness of AZ91 foam

AZ91 foam AZ91 6 mm 4.3 mm – +45%

AlMg3 2.7 mm +5%

St12 0.9 mm +7%

The upper mentioned compression and bending behaviour is valid for AZ91 as well as for AM60. In opposition to AZ91 foam AM60 foam has a higher deflection at the same force (Fig. 8). This is similar to the behaviour of the compact material. So the variation of the alloy offers the possibility to adjust the deflection. Figure 8. Three point bending test: AM60 foam has a higher deflection than AZ91 foam (density = 1.36 g/cm³).

3.4

Elastic modulus and damping properties

The elastic modulus decreases with decreasing density (Table 2). On the other hand the AZ91 integral foam has a three times better damping capacity than compact AZ91 material.

359 Table 2. Elastic modulus and damping properties AZ91 – AZ91 foam.

Density Elastic modulus Damping (logarithmic decrement)

4

AZ91 1.76 g/cm³ 38.9 GPa 4.8 x 10-4

AZ91 foam 1.34 g/cm³ 22.5 GPa 1.51 x 10-3

Summary

For the first time magnesium integral foams were produced with a special casting technique which is suitable for mass production in a cost effective way. The integral foams made out of AZ91 and AM60 have a dense skin and a foamed core. The density profile is graded and mostly symmetric with the lowest density in the core. Thin sandwich panels can be produced showing excellent mechanical properties, especially high bending stiffness at low weight, as well as high damping and energy absorbing capacity. Possible applications may be stiff light-weight panels for transportation systems, energy absorbers or housings with good vibration damping.

Acknowledgement The authors want to thank the Deutsche Forschungsgemeinschaft for financial support within the Sonderforschungsbereich 396 “Robuste verkürzte Prozessketten für flächige Leichtbauteile”. Special thanks are addressed to Andreas Lohmüller and Martin Scharrer from Neue Materialien Fürth GmbH for technical support during foam production.

References [1] C. Körner, R. F. Singer, Advanced Engineering Materials 2000, 2, 159-165. [2] M. Hartmann, K. Reindel, R. F. Singer, in Porous and Cellular Materials for Structural Applications (Ed.: D. S. Schwartz), Symposium Proceedings, Volume 521, MRS Spring Meeting, 13-17 April, 1998, Materials Research Society 1998, 211-216. [3] C. Park, S. R. Nutt, in Porous and Cellular Materials for Structural Applications (Ed.: D. S. Schwartz), Symposium Proceedings, Volume 521, MRS Spring Meeting, 13-17 April, 1998, Materials Research Society 1998, 315-320. [4] Y. Yamada, C. Wen, K. Shimojima, M. Mabuchi, M. Nakamura, T. Asahina, T. Aizawa, K. Higashi, in Magnesium Alloys and their Applications (Ed.: K. U. Kainer), Wiley-VCH, Weinheim, 2000, 645-650. [5] F.-W. Bach, D. Bormann, P.Wilk, in MetFoam 2003, Proceedings of the 3rd International Conference on Cellular Metals and Metal Foaming Technology (Ed.: J. Banhart), Berlin, 23-25 June, 2003, MIT Verlag, to be published. [6] S. Semerdjiev, Introduction to Structural Foam, Society of Plastics Engineers, Manchester, USA, 1982 [7] C. Körner, M. Hirschmann, M. Lamm, R. F. Singer, in MetFoam 2003, Proceedings of the 3rd International Conference on Cellular Metals and Metal Foaming Technology (Ed.: J. Banhart), Berlin, 23-25 June, 2003, MIT Verlag, to be published. [8] O. B. Olurin, M. Arnold, C. Körner, R. F. Singer, Materials Science and Engineering A 2002, 328, 334-343.

360

Creep of Reinforced and Unreinforced QE22 Magnesium Alloy M. Pahutova1, V. Sklenicka1, K. Kucharova1, M. Svoboda1, K.U. Kainer2 1 2

Institute of Physics of Materials, Academy of Sciences of the Czech Republic, CZ-616 62 Brno, Czech Republic Institute for Materials Research, Center for Magnesium Technology, GKSS Research Center, D-21502 Geesthacht, Germany

1

Introduction

The creep resistance of the most common magnesium AZ91 solid solution alloy (Mg9 wt.% Al-1 wt.% Zn-0.3 wt.% Mn) is rather limited at elevated temperatures. However, a very recent experimental study of the creep behavior of an AZ91 alloy reinforced with 20 vol.%Al2O3 short fibers and an unreinforced AZ91 matrix alloy has shown that the use of alumina short fibers is very effective in considerably improving the creep properties of the squeeze-cast AZ91 monolithic alloy [1,2,3]. It is now well established [4] that the creep of metal matrix composites is controlled by the rate of flow within the matrix alloys. Therefore, the creep resistant precipitate strengthened magnesium QE22 alloy was selected for further attempt to clarify the creep strengthening mechanisms in short fiber reinforced magnesium alloys. This work reports the experimental results obtained in an investigation of the high temperature creep properties of a fiber-reinforced QE22 matrix composite and its matrix alloy.

2

Experimental Materials and Procedures

Short-fiber reinforced and unreinforced blocks of high strength silver-containing alloy QE22 (Mg-2.5 wt.% Ag-2.0 wt.% Nd rich rare earths-0.6 wt.% Zr) were produced by squeeze casting. The fibre preform consisted of planar randomly distributed δ-alumina short fibres (Saffil fibres ~ 3μm in diameter with varying lengths up to an estimated maximum of ~ 150 μm). The final fibre fraction after squeeze casting in the composite was about 20 vol.%. For convenience, the composite is henceforth designated QE22 + Saffil. Unreinforced QE22 matrix alloy and its composite were subjected to a T6 heat treatment. Creep specimens were machined from the blocks so that the longitudinal specimen axes were parallel to the plane in which the long axes of the fibres were preferentially situated for the squeeze-cast composites. Tensile creep tests were carried out at temperatures from 423 to 523 K and at the applied stresses ranged from 30 to 200 MPa. Metallographic and fractographic inspections were conducted after creep testing.

3

Experimental Results

3.1

Creep behavior in the unreinforced and the reinforced QE22 alloy

Representative creep data of the QE22 alloy and the QE22 + Saffil composite at 423, 473 and 523 K are shown in Figs. 1 and 2, where the minimum creep rate ε&m is plotted

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Figure 1. Minimum creep rate versus stress for the QE22 alloy.

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Figure 2. Minimum creep rate versus stress for the QE22 + Saffil composite.

against the applied stress σ on a logarithmic scale. As demonstrated by the figures, significant differences were found in the creep behavior of the composite when compared to its matrix alloy. While the values for the apparent stress exponent n of the creep rate ε&m (n = ∂ ln ε& / ∂ ln σ )T for the unreinforced alloy decrease slightly with decreasing applied stress and increasing temperature, no curvatures indicating such dependences in the plots for the composite were found. The experimental data show that the creep behavior of the composite is similar to that of the alloy regard to the high stress exponents which suggests that the creep behavior of both materials may be controlled by precipitation strengthening due to the interaction of mobile dislocations. However, the values of n for the composite are close to ∼ 40, by comparison with the alloy exhibiting the values of n only 10-20. The higher values of n for the composite may be explained in terms of a threshold stress for creep σ0, [5]. In such case the creep behavior of the composite is not driven by the total applied stress, σ, but rather by an effective stress, σe = σ - σ0. Nevertheless, the experimental data from this work describe the creep behavior of the composite only over three orders of magnitude of creep rate or less. Thus, it is difficult to establish whether the stress exponent is genuinely constant, or increases with decreasing the applied stress indicating an existence of the threshold stress. However, it is nearly impossible from the technical reasons to perform additional tests at very low stresses. The QE22 + Saffil composite exhibits superior creep resistance compared to its monolithic alloy at a temperature of 523 K (Fig.3). This difference decreases with increasing stress and decreasing temperature and there is a tendency for the reinforcement to have no significant effect on the creep strength at a temperature of 423 K. The presence of a reinforcement leads to a substantial decrease in the overall ductility of matrix alloy. Thus, the values of the strain to fracture εf in the composite were only ∼ 1 – 2% and these values were essentially independent of stress and temperature. By contrast, the strains to fracture in the monolithic alloy were markedly higher; typically up to ∼ 30%. Experimentally determined times to fracture and strains to fracture correlate well with minimum creep rates through the Monkman–Grant relationship, ε& t f ≅ ε f ≅ const , (Fig. 4); t f is approximately inversely proportional to ε& . Note, however, that ε f itself can be strongly influenced by the mechanisms of the fracture process which could explain a shift of the composite dependence towards shorter times to fracture at the same minimum creep rates (Fig. 4). Further, Fig. 4 suggests that the creep behavior of both the composite and the matrix alloy are governed by the same deformation mechanisms. This

362 -3

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. MINIMUM CREEP RATE ε m [s-1]

Figure 3. Minimum creep rate versus stress for the QE22 alloy and the composite.

Figure 4. Dependences of time to fracture on minimum creep rate (Monkman-Grant relationship) for the QE22 alloy its composite.

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QE 22

TIME TO FRACTURE tf [s]

TIME TO FRACTURE tf [s]

suggestion is supported and perhaps more clearly illustrated in Figs. 5 and 6, where for both materials, the double logarithmic plots of the time to fracture tf, as a function of applied stress are shown. The experimentally determined values of the apparent stress exponent of the time to fracture, m, m = (∂ ln t f / ∂ ln σ )T of both the composite and the alloy are very close to the inherent values of n. The difference in the lifetimes of the reinforced and unreinforced material depends on the creep loading conditions in a similar way to that of the minimum creep rates.

423 K 473 K 523 K 100

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STRESS σ [MPa]

Figure 5. Stress dependences of time to fracture for the QE22 alloy.

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Figure 6. Stress dependences of time to fracture for the QE22 + Saffil composite.

363 3.2

Microstructural observations

Detailed microstructural investigations of the squeeze-cast QE22 monolithic alloy and the QE 22-20 vol. pct Al2O3(f) composite were reported recently by Kiehn et al. [6]. Therefore, special attention in this work was given to a study of the microstructural changes arising from creep exposure in the monolithic QE22 alloy and the QE22 + Saffil composite. In the composite, the fibers act as nucleation centers in the precipitation process promoting precipitation of Al2Nd, Mg(Ag)12Nd, and Mg3Ag phases; their morphologies can be clearly seen in Fig. 7(a) for the composite in the state prior to creep. During the creep exposure, the Mg(Ag)12Nd particles at fibers grew slightly (Fig. 7(b)) and changed their composition; the Nd/Ag ratio (in at. pct) of the initial state ≈ 0.8 changed close to 0.5 for the crept specimen (473 K, 90 MPa, and 6512 hours). The Al content in the matrix is enhanced due to the decomposition of an alumina preform binder. The evolution of the particle population inside the grains leads to the formation of Al2Nd-like cubic particles in creep at 423 K. At temperatures above 423 K, these particles are replaced by hexagonal βphase or tetragonal Mg(Ag)12Nd precipitates, which nucleate often on dislocations, as shown in Fig. 7(c). The precipitation process in the grain interiors of the unreinforced alloys is different and involves only a change of the morphological features of the tetragonal semicoherent Mg12Nd particles existing in the alloy after the T6 heat treatment. It should be noted that the preferential formation of second-phase particles along the fibermatrix interface reduces the concentration of alloying elements in the matrix so that there is a consequent decrease in the density of obstacles opposing dislocation motion.

Figure 7. TEM micrographs of QE22 + Saffil composite (a) in initial state and (b) and (c) after creep at 473 K, 90 MPa and 6512 h; showing precipitates at fiber surface and a detail of intergranular precipitates.

3.3

Fiber breakage and interfacial debonding

Creep behavior and creep resistance may be substantially influenced by the development of creep damage and fracture processes [7]. Fractographic investigation failed to reveal substantial fiber cracking and breakage during creep. Fiber breakage was restricted to a region very near to the fracture path suggesting that it occurs by the propagation of the main crack during the final stage of the creep fracture process. By contrast, an intensive debonding of the matrix/fiber interfaces was revealed at the creep fracture surfaces of the QE + Saffil composite (Fig.8). The extent of interfacial debonding was markedly increasing with decreasing creep testing temperature.

364

Figure 8. SEM micrographs of QE22 + Saffil composite showing the creep fracture surfaces (a) at 523 K, 75 MPa, and (b) at 473 K, 120 MPa.

4

Discussion

Creep strengthening of the magnesium composites may occur by either direct or indirect mechanisms. Direct strengthening is due to a load transfer from the matrix to the reinforcements. Thus, load transfer is accompanied by a redistribution of stresses in the matrix and this reduces the effective stress for creep. Indirect strengthening may be caused by the microstructural changes in the matrix of the composite which retard and/or inhibit dislocation motion and lead to a threshold stress for creep that increases the creep resistance. The QE22 + Saffil composite exhibits superior creep resistance compared to its monolithic alloy at a temperature of 523 K; this difference decreases with increasing stress and decreasing temperature. The reinforcement seems to have no significant effect on the creep strength at a temperature of 423 K. Such behavior may be explained by a temperature-dependent load transfer. In the presence of load transfer, the creep data may be successfully reconciled by putting the ratios of the minimum creep rates of the composite and the rates of the matrix alloy at the same creep loading conditions equal to a factor given n by (1-α) , where α is a load transfer coefficient having values lying within the range from 0 (no load transfer) to 1 (full-load transfer) [8]. Unfortunately, values of a load transfer coefficient α for the QE22 composite can be obtained due to insufficient creep data inferred from this study. Further, the data at very low stresses are too limited to use the procedure for evaluation of a threshold stress [5]. Observed differences in the minimum creep rates and the lifetimes of the composite and unreinforced QE22 alloy imply an existence of a critical weakening of the fiber/matrix interface. Following the results of fractographic observations an extensive debonding occurs at temperature 423 K, while a rather limited interfacial debonding was revealed at 523 K (Fig.8). Thus it can be expected that no direct strengthening due to load transfer from matrix to the fibers occurs at 423 K. In fact, in the limit of zero interface strength, where no stress can be transmitted to the fibers, the composite may be weaker than the matrix alone (Fig. 3). These results suggest that direct composite strengthening controls the creep behavior of the short-fiber QE22 composite. Indirect composite strengthening may be caused by an additional temperature-dependent precipitation processes in the composite due to the presence of the reinforcement, which leads to an attractive interaction between dislocations and incoherent second-phase particles.

365

5

Conclusions

A comparison between the creep characteristics of the high strength QE22 alloy reinforced with 20 vol.%Al2O3 (Saffil) short fibers and an unreinforced QE22 matrix alloy under comparable experimental conditions shows that creep strengthening in the composite arises mainly from effective load transfer provided that no extensive debonding of the fiber/matrix interfaces and substantial breakage of fibres during creep exposure occur. The creep behavior and the creep resistance of the composite may be substantially influenced by the creep loading conditions. The creep resistance of the composite was significantly improved by comparison with the unreinforced alloy at a temperature of 523 K. By contrast, no beneficial effect of the reinforcement on the creep rate and thus marked improvement of the creep resistance was found for the composite at 423 K. Moreover, the creep resistance of the composite is dependent on applied stress; it decreases with increasing applied stress and there is a tendency for the reinforcement to have no significant effect on the creep resistance at the highest experimental stresses.

6

Acknowledgement

Financial support for this work was provided by the Grant Agency of the Czech Republic under Grant No. 106/03/0901.

7

References

[1] V. Sklenicka, M. Pahutova, K. Kucharova, M. Svoboda, T.G. Langdon, Key Eng. Mater. 2000, 171-174, 593-600. [2] V. Sklenicka, M. Pahutova, K. Kucharova, M. Svoboda, T.G. Langdon, Metall. Mater. Trans., 2002, 33A, 883-889. [3] M. Pahutova, V. Sklenicka, K. Kucharova, M. Svoboda, Int. J. Mater & Product Technology, 2003, 18, 116-140. [4] F.A. Mohamed, in Creep Behaviour of Advanced Materials for the 21st Century (Eds. R.J. Mishra, A.K. Mukherjee, K.L. Murty), TMS, Warrendale, USA, 1999, 83-97. [5] V. Sklenicka, M. Pahutova, K. Kucharova, M. Svoboda, T.G. Langdon in Magnesium Alloys in their Applications (Ed.: K.U. Kainer), DGM, Wiley-VCH, Weinham, 2000, 246-251. [6] J. Kiehn, B. Smola, P. Vostry, I. Stulikova, K.U. Kainer, Phys. Status Solidi (a) 1997, 164, 709-723. [7] M. Pahutova, J. Brezina, K. Kucharova, V. Sklenicka, T.G. Langdon, Mater. Lett. 1999, 39, 179-183. [8] K.T. Park, E.J. Lavernia, F.A. Mohamed, Acta Metall. Mater. 1994, 42, 667-678.

366

Thermal Exposure Behaviour of Aluminum Borate Whisker Reinforced Magnesium Matrix Composites M. Y. Zheng*, K. Wu*, M. Liang*, S. Kamado**, Y. Kojima** * School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, P. R. China ** Department of Mechanical Engineering, Nagaoka University of Technology, Nagaoka 940-2188, Japan

1

Introduction

One way to develop high cost-effective magnesium matrix composites is to employ high cost-effective reinforcements. In the recent years, high cost-effective aluminum borate whisker has attracted great interest as an alternative reinforcement of SiC whisker. Its price is about 5-10% of that of SiC whisker, even lower than that of alumina short fiber. Furthermore, it has high strength, high Young’s modulus and low coefficient of thermal expansion which are competitive with those of SiC whisker. During the last decade, aluminum matrix composites reinforced with aluminum borate whisker have been extensively studied [1,2], and some Al18B4O33w/Al composite components have been applied practically in automobile engines. Recently, aluminum borate whisker reinforced magnesium matrix composites have also been developed [3-5], which exhibit almost equivalent mechanical properties to SiCw/Mg composites, and lower coefficient of thermal expansion [6]. The high cost-effective aluminum borate whisker reinforced magnesium matrix composites give a good possibility for the wide applications of magnesium matrix composites. However, the Al18B4O33 whisker surface is unstable [2], which reacts with magnesium easily during fabrication and subsequent heat treatment. For example, in Al18B4O33 whisker reinforced Mg-containing Al alloys, Al18B4O33 whisker reacts with magnesium including in aluminum alloys at about 520 °C, resulting in degradation of the composites properties [2]. Therefore, a major problem encountered in the development of aluminum borate whisker reinforced magnesium matrix composites may be the reaction between whisker and magnesium alloy matrix during fabrication, secondary process and subsequent service. However, few researches have reported on the influence of thermal exposure on the interface and properties of the Al18B4O33 whisker reinforced magnesium matrix composites. The aim of the study is to investigate the thermal exposure behaviour of the Al18B4O33w/AZ91 composite, with particular emphasis on the thermal stability of the interface between aluminum borate whisker and AZ91matrix alloy, and the effect of thermal exposure on the tensile properties of the composite.

2

Materials and experiment methods

The composites used in this investigation were fabricated by squeeze casting. The matrix alloy was commercial heat-treatable AZ91 magnesium alloy (8.5-9.5%Al, 0.45-0.90%Zn, 0.15-0.30%Mn, 0.20%Si, 0.01%Ni, balance Mg), and the reinforcement was aluminum

367 borate whisker (Al18B4O33w, Grade M12, Shikoku Chemical Corporation, Japan) with a diameter of 0.5-1.0 μm and length of 10-30 μm. The Al18B4O33 whiskers were fabricated into a preform by wet forming method without any binder. The preforms were dried in air at room temperature for 3 days, then 100 °C for 24 h. To increase the strength of the preforms, the preforms were calcined at 1300 °C for 3 hours in air. The perform was subsequently infiltrated by molten AZ91 magnesium alloy under a high pressure in a squeeze casting press. Prior to infiltration, the mold and the preform were preheated at 500 °C, and AZ91 magnesium melt superheated to 800 °C was poured over the perform. The pressure exerted by the punch on the molten metal was gradually increased to a pre-determined level of 100 MPa in order to avoid deformation of the preform, and kept at this level for 3 minutes until the molten metal was completely solidified. The volume fraction of the whiskers in the composite was about 30%. 0 After fabrication, some samples of the composite were heat treated at 400 C for different periods up to 50 h, and at different temperature up to 600 °C for 10 h, respectively. At the exposure temperature of 500 °C and 600 °C, the specimens were heat treated in evacuated glass tubes. The thin foils for TEM observation were prepared by mechanical thinning, followed by Ar ion-beam thinning. The interfaces between aluminum borate whisker and magnesium matrix were examined by Philips TECNAI 20 and CM12 transmission electron microscopes (TEM). The tensile tests were performed in Instron 1186 and Shimazu AG-1 testing machines at a crosshead speed of 0.5 mm/min. The dog bone-shaped tensile samples of the as-cast and thermal exposed composites had a gauge length of 15 mm and a thickness of 2 mm. At least three samples were tested in each condition. Tensile fracture surfaces of the specimens were characterized by Hitatch S-570 and JSM-5600 scanning electron microscopes (SEM).

3

Results and discussion

3.1

The effect of thermal exposure on the Interfacial reaction in the Al18B4O33w/AZ91 composite

Fig.1 shows typical TEM micrographs of the interface between Al18B4O33 whisker and AZ91 in the as-cast Al18B4O33w/AZ91 magnesium matrix composite. The interfacial reaction products are clearly seen at the interface, which forms a uniform and narrow interfacial reaction layer decorated at the whisker surface. The thickness of the reaction layer is about 10-20 nm. The interfacial reaction products in the composite were identified as MgO by the selected area diffraction pattern (SADP) analysis [7]. The MgO may be formed by chemical reaction between melt magnesium and aluminum borate whisker, according to the following chemical equations: Al18B4O33 + 33Mg → 33MgO +18Al + 4B, Al and B are dissolved into the matrix during the formation of MgO. In the oxide fiber reinforced magnesium matrix composites, MgO was usually observed at the interface by the interfacial reaction between oxide fiber and molten magnesium matrix [8]. During squeeze casting of the Al18B4O33w/AZ91 composite, the infiltration process is very fast, and the total solidification time is less than 1 minute. The thin uniform MgO reaction layer composed of numerous very-fine-grained MgO particles formed during

368 such short period indicates that both the nucleation rate and density of nucleation sites of MgO on the Al18B4O33 whisker surfaces are very high when the molten magnesium comes into contact with Al18B4O33 whiskers. This is supported by the observations of Lloyd et al., who have found that the nucleation rate of interfacial reaction products is higher with higher Mg content in Al2O3 particle reinforced Al-Mg alloys with different Mg content [9]. Therefore, very high nucleation rate of MgO can be expected at the Al18B4O33 whisker surface in the present composite with AZ91 containing Mg of 90% wt pct. as matrix alloy.

Figure 1. TEM micrographs of the Al18B4O33w-AZ91 interface in as-cast Al18B4O33w/AZ91 composite. (a) lower magnification and (b) higher magnification.

Fig.2 shows the interfacial microstructure of the composite after exposure at 400 °C for different periods. Up to 50 h at 400 °C, there is hardly any change in the morphology and thickness of the interfacial reaction layer compared with that in the as-cast composite. It suggests that the interface is very stable below 400 °C.

Figure 2. TEM micrographs of the interface in Al18B4O33w/AZ91composite after exposure at 400 °C for (a)18 h and (b) 50 h.

369 Fig. 3 shows the interfacial structure of the composite after exposure at different temperature for 10 h. After thermal exposure at 500 °C, as shown in Fig. 3a), interfacial reaction rate increases, leading to the formation of slightly thicker reaction layer and some slightly larger MgO particles at the matrix-reaction layer interface. After exposure at 600 °C for 10 h, as shown in Fig. 3(b), interfacial reaction rate increases obviously. The thickness of the MgO interfacial reaction layer increases to 100–200 nm. The distribution of the size of interfacial reaction products MgO in the reaction layer is uneven, larger MgO particles are formed closer to the matrix-reaction layer interface with well developed facets.

Figure 3. TEM micrographs of the interface in Al18B4O33w/AZ91composite after exposure at (a) 500 °C and (b) 600 °C for 10 h.

The uniform interfacial reaction layer formed during squeeze casting can act as a good barrier to prevent further reaction between the aluminum borate whisker and magnesium alloy matrix during subsequent thermal exposure. Only after thermal exposure above the liquidus temperature (595 °C) for a long time, the interfacial reaction rate increases obviously. The obvious interfacial reaction after exposure above the liquidus temperature indicates that the MgO reaction layer formed during squeeze casting is not a perfectly dense continuous layer, because such perfectly dense MgO layer will not grow even after exposure in the molten matrix alloy for several hundred hours, as calculated by Hallstedt [10] using reaction kinetics. During exposure at 600 °C, the molten magnesium seeps through the channels in the MgO layer formed during squeeze casting, and reacts with the whisker further. When the MgO particles closer to the reaction layer-matrix interface grow larger enough to close off the spaces between these particles, thus block the way of any metal penetration, the seepage process will stop. As the seepage mechanism [8] has been shut off by the larger MgO particles closer to the matrix-MgO layer interface, the small MgO particles closer to the MgO layer-whisker interface do not have as much time to grow as those closer to the matrix-MgO layer interface, which results to the increase of MgO particle size with distance from the MgO layer-whisker interface. However, it should be noted that the interfacial reaction rate is not so great even after exposure at 600 °C for 10 h. The uniform, dense and fine MgO particles in the reaction layer formed during squeeze casting make the diffusion channels between the fine MgO particles very narrow, reducing the diffusion of elements, and contribute to the low interfacial reaction rate.

370 In the as-cast aluminum borate whisker reinforced Al-Mg aluminum alloy composites[2], the Al18B4O33 whiskers are heavily damaged by the serious reaction between Al18B4O33 whiskers and Mg in the Al-Mg alloys during the squeeze casting and subsequent heat treatment, leading to deterioration of reinforcing efficiency of Al18B4O33 whisker. Therefore, in order to prohibit the reaction between the aluminum borate whisker and the liquid aluminum-magnesium alloys, surface treatment techniques of Al18B4O33 whisker have to be developed to establish appropriate barrier layer. In the aluminum borate whisker reinforced AZ91 magnesium matrix composites, the uniform MgO reaction layer formed at the interface during squeeze casting can effectively prevent the further reaction between Al18B4O33 whisker and magnesium matrix, and the surface treatment of Al18B4O33 whisker seems unnecessary. The interface between aluminum borate whisker and magnesium alloy can be thought to be more stable than the interface between aluminum borate whisker and Al-Mg alloy. The effect of thermal exposure on the tensile properties of the Al18B4O33w/AZ91 composite

3.2

The aluminum borate whisker is very effective in improving the mechanical properties of AZ91 magnesium alloy. The effect of thermal exposure period and temperature on the tensile properties of the Al18B4O33w/AZ91 composite are shown in Fig. 4 and Fig. 5, respectively. There is almost no deleterious effect on the mechanical properties of the composite after thermal exposure below the solidus temperature of AZ91 alloy for a long time. These correlate with a very thin and even MgO reaction layer at the interface, which is generally considered as beneficial because it enhances the interfacial bonding between matrix and reinforcement [11]. After exposure at 500 °C for 10 h, the tensile properties of the composite decrease significantly, although the interfacial reaction layer only thickens slightly. This property degradation may be mainly originated from the porous formed in the specimen, which resulted from the evaporation of magnesium during exposure at high temperature in a vacuum. 450

500 0

Exposure at 400 C

350 300 250 200 150 100 50 0

Exposure for 10h

450

Ultimate tensile strength, MPa

Ultimate tensile stress, MPa

400

400 350 300 250 200 150 100 50

0

10

20

30

40

50

Exposure time, h

Figure 4. Effect of thermal exposure time at 400 °C on ultimate tensile stress of Al18B4O33w/AZ91 composite.

0

0

100

200

300

400

500

o

Exposure temperature, C

Figure 5. Effect of thermal exposure temperature for 10 hours on ultimate tensile stress of Al18B4O33w/AZ91 composite.

The tensile fractographs of the Al18B4O33w/AZ91 composite under different conditions are shown in Fig. 6. The overall fracture surfaces of the composites are generally unaffected by thermal exposure. The interfacial bond in all the composites keeps strong although the interfacial reaction occurs.

371

4

Conclusions

Figure 6. SEM fractographs of the Al18B4O33w/AZ91 composite. (a) as-cast and (b) after exposure at 400 °C for 50 h.

A uniform and dense MgO interfacial reaction layer with a thickness of 10-20 nm is formed at the matrix-whisker interface in the Al18B4O33w/AZ91 magnesium matrix composite during squeeze casting, which resulted form the interfacial reaction between Al18B4O33 whisker and Mg matrix. The MgO reaction layer formed during squeeze casting can act as a good barrier to prevent the further interfacial reaction during subsequent thermal exposure of the composite. The interfacial reaction began to increase obviously only after the composite was held above the liquidus temperature of AZ91 for a long periods. Thermal exposure below the solidus temperature of AZ91 matrix alloy for a long periods has almost no deleterious effect on the mechanical properties of the Al18B4O33w/AZ91 composite, indicating that the composite can be subjected to long- term service at elevated temperature.

5 [1] [2]

References

K. Suganuma, T. Fujita, N. Suzuki, K. Niihara, J. Mater. Sci. Lett. 1990, 9, 633-635. J. Pan, G. Sasaki, L. J. Yao, M. Yoshida, H. Fukunaga, Mater. Sci. Tech. 1999, 15, 1044-1048. [3] S. Kamado, T. Shinkawa, T. Wada, Y. Kojima, J. Japan Inst. Light Metals. 1995, 46, 71-76. [4] G. Sasaki, M. Yoshida, J. Pan, N. Fuyama, T. Fujii and H. Fukunaga, Mater. Sci. Forum. 2000, 350-351, 215-220. [5] M. Y. Zheng, K. Wu, H. C. Liang, S. Kamado, Y. Kojima, Mater. Letter. 2002, 57, 558-564. [6] M. Y. Zheng, K. Wu, C. K. Yao, Mater. Sci. Eng. A. 2001, A318, 50-56. [7] M. Y. Zheng, K. Wu, M. Liang, S. Kamado, Y. Kojima, J. Mater. Sci. Lett. 2003 (In press). [8] M. Pfeifer, J. M. Rigsbee and K. K. Chawla, J. Mater. Sci. 1990, 25, 1563-1567. [9] D. J. Lloyd, I. Jin and G. C. weatherly, Scr. Metall. 1994, 31, 393-397. [10] B. Hallstedt, Z. K. Liu, J. Agren, Mater. Sci. Eng. A.1990, 129, 135-145. [11] S. Ochiai, Y. Murakami, J. Mater. Sci. 1978, 14, 831-840.

372

Creep Properties of Mg Dispersion-strengthened by Graphite Hans Ferkel Institut für Werkstoffkunde und Werkstofftechnik, Technische Universität Clausthal, Agricolastr. 6, 38678 Clausthal-Zellerfeld, Germany.

Abstract Mg was dispersion-strengthened by graphite via powder metallurgy. The material was produced by ball milling of Mg micropowder of a median particle diameter of 40 μm with 3 Vol.% of graphite powder with a median particle diameter between 1 and 2 μm. After 8 h of milling the product was consolidated by hot extrusion. Structure analysis reveals that a submicrocrystalline structure developed during milling. Tensile tests show that the material is brittle even up to 300 °C and therefore, most mechanical tests were carried out under compression. Under those condition the reinforced material shows flow stresses of 270 MPa at ambient temperatures, 170 MPa at 150 °C, and 125 MPa at 300 °C whereas Mg processed under the same condition without graphite addition reveals significantly lower yield stresses. The dispersion-strengthened Mg shows a marked increase in the creep resistance: at 200 °C and a stress σc of 100 MPa, the secondary creep rate εs is in the lower 10-9 s-1 range and at 300 °C and σc of 80 MPa, εs values of up to 1 x 10-8 s-1 were measured. The results are discussed.

1

Introduction

Mg, its alloys and Mg matrix composites received increasing attention within the last decade due to their high potential as light-weight materials in automotive and aerospace applications [1-4]. A major disadvantage of Mg is its limited strength especially at elevated temperatures and therefore, new precipitation hardened Mg alloys were developed. Unfortunately, these alloys often show deterioration of the corrosion properties [2]. In order to achieve acceptable features, efforts are spent to develop Mg matrix composites. Furthermore, dispersion-strengthening of Mg by a low volume fraction (≤3 Vol.%) of a second inert nanoscaled well dispersed ceramic phase is a promising approach to obtain materials of higher strength and improved creep resistance [5-7]. In this work preliminary results of Mg stengthened by fine graphite dispersed via powder metallurgy in a pure Mg matrix are presented.

2

Experimental

The Mg-micropowder employed was produced by gas atomisation of a HP magnesium melt with argon containing 1 % oxygen for powder passivation. The powder had a median

373 particle diameter of about 40 μm. The graphite powder used had a median particle size between 1 and 2 μm. The Mg-powder was mixed for 8 hours with 3 Vol.% of the graphite powder in a asymmetrically moved mixer followed by milling for 8 h at 200 rpm in a planetary ball mill (Retsch, PM400) in a sealed Argon atmosphere. The milling vessel of 500 ml volume was made of corundum and the milling balls having a diameter of 11 mm were made of hardened stainless steel (100Cr6). The weight ratio of ball-to-powder was 10:1. The composites were encapsulated in an evacuated magnesium container (70 mm in diameter), degased at about 350 °C and extruded by a preheated (350 °C) 400 t horizontal extrusion press (outlet 14 mm). For comparison also Mg powder (no additives) was milled and consolidated in the same way. In order to study the materials regarding contamination they were analysed in an optical spark analyser (Spectrolab, Spectro Analytical Instruments). Tensile tests of the as-prepared materials were performed with a PC controlled universal -4 -1 testing machine at strain rate dε/dt = 3.3 × 10 s at different temperatures in air. Creep tests were carried out under compression at constant stresses at 200 °C and 300 °C in air. All specimens tested were machined from the extruded bars with their load axis parallel to the extrusion direction. The microstructure of the materials was investigated by light microscopy.

3

Results and Discussion

Figure 1 shows light microscopic pictures of the transversal section of the pure Mg and the C/Mg composites. As can be seen, both materials developed a very fine-grained structure after processing with the structure of the re-inforced material being the finest (right part of Fig. 1). All specimens show elongated grains parallel to the extrusion direction (not shown). Optical spark analysis revealed that the materials were not contaminated by impurities as iron from the extruder tools or from the milling balls and Al from the milling vessel.

Figure 1. Light microscopic pictures of the milled and extruded material. Mg (left part) and Mg strengthened by 3 Vol.% of graphite (right part).

Tensile tests were carried out under compression and the results are summarised in Fig. 2. At all testing temperatures the re-inforced material exhibits yield stress σ0,2 and compressive strength σB values superior to pure Mg. It should be mentioned that tension tests were also carried out. However, the re-inforced material is brittle and only at 300 °C it shows values of σ0,2 and UTS exceeding those of pure Mg.

374 500

σ 0.2

σB

RT

450

150°C

Mg

400

Mg + 3 Vol.% C

Mg

Mg + 3 Vol.% C

σ in MPa

350

300°C

350 300 250

270 245 225 198

200

170

150 100

125 105 84

135

120 88

50 0

Figure 2. Yield stress σ0,2 and compressive strength σB of Mg and Mg + 3 Vol.% C at various temperatures.

Figures 3-5 present creep results. Only the sections of secondary creep are shown. As can be seen by comparing Figs. 3 and 4, the creep rate of Mg was significantly improved by the incorporation of graphite. The gain in strength could be explained in the context of a low dislocation density and/or pinning of dislocations and hindered Mg grain sliding by the dispersed graphite phase. However more structure investigations are planned to elucidate the situation. The secondary (steady-state) creep rate εs is often described by standard power low relationship dεσ/dt = A ⋅ σn ⋅ D-m ⋅ exp[-Qc/RT] where A is a constant of proportionality, n is the stress exponent, σ the applied stress, D the (matrix) grain size, m ≥ 1, Qc the activation energy for the rate controlling deformation process, R is the gas constant and T is the absolute temperature. In the case of dislocation controlled creep of dispersionstrengthened materials, the standard creep rate equation is often refined by introducing a threshold stress necessary to overcome a dispersoid by a dislocation [e.g. 8]. Rösler and Arzt developed a detachment-controlled creep rate model in which a thermal activated dislocation detachment on the departure side of the dispersoid controls the creep rate depending on the relaxation of the dislocation strain field in the vicinity of the particle/matrix interface which can result in an attractive interaction between dislocation and dispersoid [9]. However these equations are less applicable for polycrystalline (fine grained) materials. Nevertheless, a plot of dεs/dt as a function of σ in a double logarithmic representation gives a first glance of the stress exponent n of the materials. Fig. 5 shows a power law representation recorded at different stresses in tests carried out at 200°C. In particular the re-inforced material shows a stress exponent larger than 10, which is also observed in Mg dispersion-strengthened by SiC nanoparticles [7]. Obviously, a power law breakdown can be observed for this material. It should also be mentioned that the secondary creep rate of milled pure Mg tested at 200 °C and 45 MPa under compressive load is more than 2 orders of magnitude smaller than Mg (produced via powder metallurgy) having an average grain size of around 40 μm. Due to the much smaller grain size, the opposite behaviour would be expected. However,

375 one has to keep in mind that the Mg-powder employed is slightly contaminated by oxygen on the particle surface. Therefore, milling of such powder will result in a dispersion of MgO in the Mg matrix. This might affect the deformation behaviour similar as other dispersoids as C or SiC but to a lower extent. In Fig. 4 (lower part) also one creep curve of the re-inforced material measured at 300 °C and 80 MPa is given. This result demonstrates the high creep resistance of the material even at relatively high temperatures. More creep tests at different temperatures and constant stresses will allow to extract the activation energy Qc for the rate controlling deformation process from a plot of ln(dεσ/dt) against (1/T) for different stress regimes. In conclusion, it was found that dispersion-strengthening of Mg by graphite via comilling followed by hot extrusion is an interesting approach to improve the creep properties of Mg. In comparison to earlier studies, where laser-generated SiC nanoparticles were dispersed in Mg via the same PM route [7], graphite compares well with SiCnanoparticles at least from an economical point of view. t in s 0,0E+00 -5,140E-02

2,0E+05

4,0E+05

6,0E+05

8,0E+05

1,0E+06

-5,160E-02

-5,180E-02

ε

-5,200E-02

-5,220E-02 Mg pure milled, 200 °C, 60 MPa -5,240E-02

Mg, 200°C, 60MPa; dε/dt = 1.7 × 10-9 s-1

-5,260E-02

-5,280E-02

t in s

0,0E+00 2,0E+05 4,0E+05 6,0E+05 8,0E+05 1,0E+06 1,2E+06 1,4E+06 1,6E+06 1,8E+06 -5,250E-02 -5,300E-02 -5,350E-02

ε

-5,400E-02 -5,450E-02 -5,500E-02 -5,550E-02

Mg, 200°C, 70MPa; dε/dt = 3.5 × 10-9 s-1

-5,600E-02 -5,650E-02

Figure 3. Creep curves (section of secondary creep) of the milled pure Mg under compression.

376 t in s

0,0E+00 1,0E+05 2,0E+05 3,0E+05 4,0E+05 5,0E+05 6,0E+05 7,0E+05 8,0E+05 -1,80E-02 -1,82E-02 -1,84E-02

ε

-1,86E-02 -1,88E-02 -1,90E-02

Mg+3Vol.%C, 200°C, 115MPa;

dε/dt = 3 × 10-9 s-1

-1,92E-02 -1,94E-02

0,0E+00 -1,85E-02

5,0E+04

1,0E+05

1,5E+05

t in s

2,0E+05

2,5E+05

3,0E+05

3,5E+05

-1,90E-02 -1,95E-02

ε

-2,00E-02 -2,05E-02 -2,10E-02 -2,15E-02 -2,20E-02

Mg+3Vol.%C, 200°C, 135MPa;

dε/dt = 1.5 × 10-8 s-1

-2,25E-02

0,0E+00 -1,05E-01

5,0E+04

1,0E+05

t [s]

1,5E+05

2,0E+05

2,5E+05

3,0E+05

-1,06E-01 -1,07E-01 -1,08E-01

ε

-1,09E-01 -1,10E-01 -1,11E-01 -1,12E-01 -1,13E-01 -1,14E-01

Mg+3Vol.%C, 300°C, 80MPa;

dε/dt = 4 × 10-8 s-1

-1,15E-01

Figure 4. Creep curves (section of secondary creep) of milled Mg reinfored by 3 Vol.% of graphite under compression. The two upper figures present two measurements taken at 200 °C and the measurement in the lower figure was taken at 300 °C.

377 1,0E-04

dε/dt in s-1

1,0E-05

1,0E-06

1,0E-07

1,0E-08

1,0E-09

1,0E-10 10

100

σ in MPa

1000

Figure 5. log(dε/dt)/logσ plot of the steady-state creep rates recorded at different compressive stresses in tests carried out at 200 °C for Mg (circles) and Mg reinforced by 3 Vol.% of graphite (squares).

References [1] B.L. Mordike, K.U. Kainer, Magnesium Alloys and their Applications,WerkstoffInformationsgesellschaft, Frankfurt, Germany (1998). [2] G. Neite, K. Kubota, K. Higashi, F. Hehmann, in: R.W. Cahn, P. Haasen, and E.J. Kramer (Eds.), VCH-Verlag, Mater. Sci. Tech. 8 (1996) 114. [3] D.J. Lloyd, Intern. Mater. Rev. 39 (1994) 1. [4] R. Oakley, R.F. Cochrane, R. Stevens, Key Eng. Mater. 104 (1995) 387, and references therein. [5] J. Naser, W. Riehemann, H. Ferkel, Mater. Sci. Eng. A 234 (1997) 467. [6] R. Unverricht, V. Peitz, W. Riehemann, H. Ferkel, in [1], p. 327. [7] H. Ferkel, B.L. Mordike, Mater. Sci. Eng, A298 (2001) 193. [8] R.W. Lund, W.D. Nix, Acta Met. 24 (1976) 469. [9] J. Rösler, E. Arzt, Acta Met. 38 (1990) 671.

378

High Strength Si-O-C Ceramic Foams for the Reinforcement of Mg-Alloys J. Zeschky1, T. Höfner1, J. Lo2, M. Scheffler1, P. Greil1 1 2

University of Erlangen-Nuremberg, Erlangen, Germany Canada Centre for Mineral and Energy Technology, Ottawa, Ontario, Canada

Abstract Magnesium alloys are gaining interest as light weight construction elements in automotive applications. In the temperature range exceeding 100 °C Mg alloys feature low strength and stiffness. The reinforcement of cast Mg alloys using rigid ceramic foams results in a significant improvement of the strength and creep resistance at elevated temperatures. Open cell polymer derived Si-O-C ceramic foams were fabricated using filler loaded poly (siloxane) by an in situ foaming process at 270 °C and subsequent pyrolysis at 1000 °C in nitrogen atmosphere. A 65 wt.-% Si/SiC blend was used as inert filler. The ceramic foams are characterised by a high interconnectivity of the foam cells and the pore diameter ranges from 0.5 mm to 2.0 mm. The compression strength exceeds 4 MPa at a fractional density of 27%. Interpenetrating Mg-ceramic-composites were manufactured by squeeze casting that offer improved compression strength and reduced creep rate in the temperature range above 100 °C.

1

Introduction

Ceramic foams offer high mechanical stability and stiffness at low density. The foam structure consists of a linked network of irregular formed open or closed cell polyhedrons. Due to its low density cellular structure ceramic foams exhibit non catastrophic localized fracture behaviour and superior thermal shock resistance that make them interesting candidate materials for conventional (metal melt filtration, catalyst carrier) as well as novel applications (porous gas burners, bio technology, reinforcement of light metals [1]). Ceramic foams can be manufactured using replica techniques [2] (dip coating of a polymer foam and subsequent burn-out of the polymer), space holder methods [3] (infiltration of salt compacts and dissolving the salt afterwards), gel casting methods [4] and gas bubble formation techniques [5]. In the recent years, silicone resins were used for the fabrication of ceramic foams [6]. Poly (silsesquioxane) was foamed using poly (urethane) as foaming agent. The foam structure could be modified by adding tensides to the polymer. The fabrication route applied in this study facilitates the fabrication of ceramic foams without adding further foaming agents [7]. A phenyl methyl poly (silsesquioxane) containing small amounts of ethoxy- and hydroxy-groups was foamed in situ. Heating in the temperature range of 220 – 300 °C gives rise for the cross linking of the polymer via condensation reactions:

379 ≡Si-OH + HO-Si≡

⎯→

≡Si-O-Si≡ + H2O↑

(1a)

(1b). ≡Si-OC2H5 + HO-Si≡ ⎯→ ≡Si-O-Si≡ + C2H5OH↑ The release of gaseous water and ethanol leads to the formation of foam cells. The viscosity of the polymer melt increases with increasing degree of curing which stabilizes the foam structure. Finally the foam thermoset is pyrolysed resulting in a Si-O-C based ceramic. The term “Si-O-C” denotes the partial crystalline silicon oxycarbide with free carbon clusters of 2-5 nm randomly distributed in an amorphous Si-O-C network [8]. Loading the polymer with reactive or inert filler particles to the polymer, reaction products are formed that may be used to tailor the mechanical, magnetical and electrical properties [9]. In this study, Si/SiC filled ceramic foams were infiltrated with Magnesium alloy AZ91. The metal/ceramic composites were characterized regarding the microstructure and the mechanical properties at elevated temperatures. Compared to AZ91 the composite exhibits a creep rate 80 times lower and a 50% improvement of the Young’s modulus and of the compression strength at 135 °C and 200 °C.

2

Experimental Procedure

A methyl phenyl poly (silsesquioxane) ([(C6H5)0.62(CH3)0.31(OR)0.07SiO1.5]n with n ~ 20, R= [–OH] and [–OC2H5], Silres H44, Wacker Chemie, Burghausen, Germany) was taken as preceramic polymer. The polymer is solid at room temperature, has a meting interval of 40 – 60 °C and a density of 1.1 g/cm3. A Si/SiC blend (Si: Silgrain 0-45 μm, Elkem, Meerbusch, Germany; SiC: F500B, H.C. Starck, Selb, Germany) was added to the polymer in a polymer: Si: SiC weight ratio of 35:12:53. Si and SiC were selected as inert fillers to reduce the shrinkage and to increase the strength and structure stability of the foam [10]. The blend was mixed in a ball mill for 2 hours. The mixture was foamed in an Aluminium tray at 270 °C for 2 hours in air using a preheated convection furnace. The thermoset foam was cut in discs (r = 35 mm, h = 10 mm). One set of foams was fired in air at 1000 °C for 4 hours, another set of foam was pyrolysed in nitrogen at 1000 °C for 4 hours and oxidized in air at 1000 °C for 12 hours. The samples were heated by 3 °C/min to 500 °C and by 3 °C /min to 1000 °C. The cooling rate was set to 3 °C /min. The strut thickness, cell diameter, connectivity density and surface/volume ratio were determined by x-ray micro computer tomography μCT (μCT 40, Scanco Medical GmbH, Bassersdorf, Switzerland). The compression strength was measured at room temperature 3 for each 10 cuboidal shaped samples (10 * 10 * 20 mm ) (Instron 4202, Instron Corp., Canton, MA, USA). The speed of the cross head was set to 0.5 mm/min. The metal melt infiltration with AZ91 (90 wt.% Mg, 9 wt.% Al, 1 wt.% Zn, traces of Mn) was carried out under a pressure of 86 MPa at 680 °C by squeeze casting in CO2 atmosphere [11]. The pressure was remained for 2 minutes until the melt solidified. For the stress-strain measurements cubic samples (7 * 7 * 7 mm3) were cut and tested at 135 °C and 200 °C at a constant strain rate of 10-4 s-1 (Instron 4202, Instron Corp., Canton, MA, USA). Creep tests were performed at pyrolysed, oxidized, and infiltrated foam samples of the same geometry at 200 °C. A constant compression stress of 100 MPa was applied. The strain was measured using an inductive position encoder. The Young’s modulus was determined by measuring the ultrasonic speed through the material (USD10, Krautkramer Branson, Marina del Rey, CA, USA). Scanning electron microscopy (Quanta 200, FEI Company, Eindhoven, Netherlands) was applied to cut and polished samples.

380

3

Results and Discussion

3.1

Ceramic foams

Figure 1 shows a 3D μCT reconstruction of a polymer derived Si-O-C ceramic foam. After pyrolysis in nitrogen and subsequent oxidation in air or after firing in air a completely open cell structure was obtained. Oxidation of the surface has shown to improve interface integrity between the metal and ceramic phase during infiltration and to prevent from debonding during cooling [10]. The foams show a porosity of 73% having a narrow strut thickness distribution with a mean thickness of 0.44 ± 0.19 mm and a mean cell diameter of 1.30 ± 0.36 mm with a small shoulder at 0.6 mm, Figure 2.

1 cm Figure 1. 3D micro computer tomography of a Si/SiC filled Si-O-C foam pyrolysed in nitrogen atmosphere at 1000 °C for 4 h and oxidized in air at 1000 °C for 12 h.

-1

Probability Density (mm )

6

Strut Thickness

5

Cell Diameter

4 3 2 1 0 0.0

0.2

0.4

0.6

0.8

1.0 1.2 Size (mm)

1.4

1.6

1.8

2.0

Figure 2. Strut thickness and cell diameter distributions of the pyrolysed Si-O-C foam.

The foams exhibit a connectivity density of 1.0 mm-3 and a surface/volume ratio of 6.5 mm-1. Foams that were pyrolysed in nitrogen and subsequently oxidized consist of an amorphous matrix surrounding the Si and SiC particles. Due to the oxidation step, the surface of the foam is covered with a thin layer of SiO2 (cristolbalite). Foams that were fired in air did not reveal any remaining carbon as proofed by EDX analysis. Due to the carbon burn-

381 out the strut matrix shows an increased porosity. The crushing strength σ* and the Young’s modulus E* of the pyrolysed foam was 3.9 MPa and 3.2 GPa, whereas the foams fired in air exhibited σ* = 3.4 MPa and E* = 2.7 GPa, respectively. The micro porosity in the struts of the foams fired in air result in a decrease of the crushing strength and of the Young’s modulus of about 15 %. 3.2

Interpenetrating phase composites

Due to the open cell structure of the foams, completely infiltrated interpenetrating phase composites (IPCs) were obtained after squeeze casting, Figure 3.

Figure 3. SEM micrograph of the Si-O-C ceramic/metal alloy interface (dark: ceramic skeleton, grey: Mg matrix, white: Mg2Si and Al12Mg17 precipitations).

It is essential for an effective reinforcement of the metal matrix to generate a strong interface bonding of the metal to the ceramic skeleton. In the case that no interface reactions occur, debonding of the phases due to the mismatch of the coefficient of thermal -6 -1 -6 -1 expansion (CTE(AZ91) = 28 ⋅ 10 K , CTE(Si-O-C foam) = 8 ⋅ 10 K ) is inevitable resulting in an IPC material with extended debonding cracking between the metal matrix and the ceramic skeleton. Hence load transfer from the matrix to the reinforcing skeleton is restricted and no reinforcing effect can be expected. Thus, formation of an interface bonding layer should be highly favourable10. From experimental observation the following reaction SiO2 (s) + 4 Mg(l) ⇒ 2 MgO (s) + Mg2Si

(ΔH = -410 kJ/mole at 680 °C)

is highly probable to occur and to give rise for a crack free interface bonding and a smooth transition of the thermal induced stresses at the interface. Precipitations of Mg2Si and Al12Mg17 were found in the metal phase and Si, SiC, MgO and small amounts of Mg2Si were detected in the strut material.

382 IPCs loaded at 135 °C and 200 °C showed an increased Young’s modulus and crushing strength as compared to the unreinforced Mg alloy AZ91. The pyrolysed/oxidized foam IPC has a Young’s modulus of 67 GPa and the oxidized foam IPC E = 55 GPa. This is an increase of 50 % and 22 %, respectively, compared to E = 45 GPa of the unreinforced Mg alloy. Figure 4 shows the true stress – true strain diagrams of AZ91 and both types of AZ91 reinforced by Si-O-C foams at 135 °C and 200 °C. The improvement is retained at elevated temperatures resulting in an elastic modulus twice as high as the unreinforced Mg alloy at 200 °C. The crushing strength of the IPC material made from the foam that was fired in air is similar to that of AZ91 at temperatures up to 135 °C. However, the decrease of the crushing strength with rising temperature is far less than for AZ91. Thus, at 200 °C the crushing strength of the IPC is 1/3 higher than of AZ91. The reinforcement effect of the IPC made from the foam which was pyrolysed in nitrogen and oxidized is more emphasized. In the whole temperature range from RT to 200 °C, the crushing strength is 50 – 60 % higher as compared to the unreinforced AZ91. a) 400

200 °C

IPC, pyrolyzed and oxidized foam 200

300

True Stress (MPa)

True Stress (MPa)

b) 250

135 °C

IPC, pyrolyzed and oxidized foam

350

250 AZ91

200 150 100

0 0.00

0.05

0.10

0.15

0.20

0.25

AZ91 100 IPC, foam fired in air 50

IPC, foam fired in air

50

150

0 0.00

0.30

0.05

0.10

0.15

0.20

0.25

0.30

True Strain

True Strain

Figure 4. True compression stress – true strain diagrams of unreinforced AZ91 and AZ91 reinforced by ceramic foams fired in air or pyrolysed in nitrogen and subsequently oxidized in air at 135 °C (a) and 200 °C (b).

The improvement of the compression strength is confirmed by results from creep tests. Figure 5 shows the true strain – time and creep rate – true strain diagrams of bulk AZ91 and ceramic foam reinforced AZ91. For the IPC a steady state creep rate of less than -7 -1 1.510 s was reached after 1 hour and remained constant until the ceramic skeleton failed at a strain of 0.02 after 25 hours. The unreinforced AZ91 showed a steady state creep rate of 8 · 10-7 s-1 at a true strain of 0.08. b)1E-4

0.20

a)

200 °C 100 MPa

AZ91

200 °C 100 MPa

True Strain

-1

Creep rate (s )

0.15

0.10 IPC, pyrolyzed and oxidized foam

AZ91

1E-5

1E-6

0.05 IPC, oxidized and pyrolyzed foam

0.00

0

10

20

30 Time (h)

40

50

60

1E-7 0.00

0.05

0.10

0.15

True Strain

Figure 5. True strain – time diagram (a) and creep rate – true strain diagram (b) of unreinforced AZ91 and AZ91 reinforced by a pyrolysed and oxidized ceramic foam under a compression load of 100 MPa at 200 °C.

383

4

Conclusion

Open cell Si-O-C-ceramic foams can be fabricated from filler loaded poly (silsesquioxane) by in situ blowing of the silicone resin in an one step manufacturing process. The ceramic foams were used for the reinforcement of Mg alloys at temperatures above 100 °C. The strength and Young’s modulus of the composites could be increased by 50% and the creep rate reduced by a factor of > 50 as compared to the unreinforced AZ91.

References [1]

L. D. Wegner and L. J. Gibson, "The mechanical behaviour of interpenetrating phase composites – I: modelling", International Journal of Mechanical Science, 42 [5] 925-942 (2000). [2] K. Schwartzwalder and A.V. Somers, Method for Making Porous Ceramic Articles, U.S. Patent, USA, 1963. [3] T. J. Fitzgerald, V. J. Michaud, and A. Mortensen, "Processing of Microcellular SiC Foams. Part II. Ceramic Foam Production", Journal of Material Science, 30 10371045 (1995). [4] Pilar Sepulveda, "Gelcasting Foams for Porous Ceramics", American Ceramic Society Bulletin, 76 [10] 61-65 (1997). [5] X. Bao, M. R. Nangrejo, and M. J. Edirisinghe, "Sythesis of Silicon Carbide Foams from Polymeric Precursors and their Blends", Journal of Material Science, 34 [11] 2495-2505 (1999). [6] P. Colombo and J. R. Hellmann, "Ceramic foams from preceramic polymers", Materials Research Innovations, 6 [5-6] 260-272 (2002). [7] T. Gambaryan-Roisman, M. Scheffler, P. Buhler, and P. Greil, "Processing of Ceramic Foam by Pyrolysis of Filler containing Phenylmethyl Polysiloxane", Ceramic Transactions, 108 121-130 (2000). [8] G. M. Renlund, S. Prochazka, and R. H. Doremus, "Silicon Oxycarbide Glasses: Part II. Structure and Properties", Journal of Materials Research, 6 [12] 2723-2730 (1991). [9] P. Greil, "Active Filler Controlled Pyrolysis of Preceramic Polymers", Journal of the American Ceramic Society, 78 835-848 (1995). [10] J. Zeschky, J. Neubauer, J. S. H. Lo, M. Scheffler, B. Kummer, and P. Greil, "Verstärkung von Magnesiumlegierungen mit Keramikschäumen aus präkeramischen Polymeren"; pp. 67-72 in 14. Symposium Verbundwerkstoffe und Werkstoffverbunde (in German). Edited by H. P. Degischer. Wiley-VCH, Vienna (Austria), 2003. [11] J. S. H. Lo and G. J. C. Carpenter, "Fabrication of SiC-reinforced AZ91D Magnesium based Composites"; pp. 688 in Eleventh International Conference on Composite Materials, 3, 1997.

384

Creep Behavior of Mg-Alloys AE42 and AZ91 Reinforced with Carbon Short Fibers S. Ataya1, S. Mielke2, E. El-Magd3 1, 3 2

Department of Materials Science - LFW, RWTH, Aachen, Germany Kolbenschmidt, Neckarsulm, Germany

1

Introduction

To realize the application potential of the light Mg-alloys, the strength at elevated temperatures and the elastic modulus of these alloys should be improved. That could be achieved by reinforcing of the Mg-alloys with a high strength component. Magnesium alloys could be reinforced with ceramic particles, continuous fibers or short fibers from Al2O3, SiC or C. With applying of reinforcement in the form of fibers or particles, particularly in the reciprocating components, a weight saving optimization should be taken into consideration that leads to selection of carbon fibers instead of ceramics. Due to the high cost of materials reinforced with continuous carbon fiber, it is meaningful to use carbon short fibers as a reinforcing material [1,2]. The reinforcing of the Mg-alloy AZ91 with short fibers [3-5] has improved the strength and creep resistance. This work aims to experimental evaluation of the creep behavior of the Mg-alloys AE42 and AZ91 reinforced with carbon short fibers (23 vol. %C) compared with the matrix alloys in terms of the load carrying capacity, creep life and creep strain. Determination of the dominant creep deformation mechanism has also a significant attention in this work. Furthermore, the anisotropic creep deformation on loading either parallel or normal to the reinforced plane due to the fibers distribution will be explored to show how much these short fibers could improve the creep resistance in a specific direction.

2

Studied Materials and Experimental Procedure

Magnesium-based composites were produced by reinforcing of magnesium alloys AZ91 (Mg-9 Al, 1 Zn, 0.5 Si, 0.13 Mn) and AE42 (Mg-4 Al, 2 RE, 0.3 Mn) with carbon short fibers PAN. These composite materials were produced by the squeeze casting technique in the company Kolben-Schmidt, Germany. Fibers were quasi-isotrpic distributed in the horizontal plane. The fibers content was about 23 vol. % with the geometry L ≈ 92 μm and D ≈ 7 μm. Cylindrical specimens with a height of H = 8 mm and a diameter of D = 6 mm were used. Creep tests were carried out on the matrix alloys and the composite materials at the temperatures 150 °C, 200 °C and 250 °C. The specimens from AE42-23 %C were machined to be tested under compressive creep by loading either parallel (Figure 1.a) or normal to the reinforced plane (Figure 1.b). While the specimens from AZ91-23%C were tested only by loading parallel to the reinforced plane (Figure 1.a).

385 Microstructure was investigated under the optical microscope as well as under SEM equipped with EDX-analyzer.

(a)

(b)

Figure 1: Arrangement of the applied load with respect to the reinforcing in the Mg-composites. (a) The load applied parallel to the reinforced plane, and (b) the load applied normal to the reinforced plane.

3

Results and Discussion

Relating to the creep behavior of matrix alloy AE42 at 200 °C (Figure 2.a,b), higher improvement was noticed in the creep resistance by loading parallel to the reinforced plane (Figure 2.e,f) than that on loading normal to the reinforced plane (Figure 2.c,d). That could be clearly noticed in the form of increasing the load carrying capacity or the life under specific

Creep strain ε, %

110

(a)

30

90

20

85

(c)

130 120

80

110 100 90

75 70 / MPa

10

AE42-23%C- parallel / 200°C

AE42-23%C- normal / 200°C

Matrix-AE42 / 200°C

40

85 / MPa

545 h

(e) 165 / MPa 160 155 150 145

140

130

0 100

dε/dt, h-1

10-1

(b) 110

10-2

90

85 80

10-3

75

(d) 100

70 / MPa 85 / MPa

0

100

165 / MPa 160 155 150

90

10-4 10-5

130 120 110

200

Time t, h

300 0

100

200

Time t, h

(f)

145

140 130

545 h

300 0

100

200

300

Time t, h

Figure 2: Compressive creep curves and creep rate against the time at 200 °C for (a,b) the matrix alloy AE42andits composite material loaded (c,d) normal and (e,f) parallel to the reinforced plane.

386 stress and the decrease of the creep strain up to fracture in case of the reinforced material especially on loading parallel to the reinforced plane. The creep behavior of the matrix alloys AE42 and AZ91 and their composite materials could be distinguished according to the double logarithmic representation of the creep stress (σ) against the minimum creep rate (ε&min ) as shown in Figure 3 according to the relation: ε&min = Bσ n exp

( ) −Q RT

dε/dtmin , s-1

where, n is the Norton-Bailey stress exponent, B is a constant, Q is activation energy for the creep deformation, R is the gas constant and T is the absolute temperature. At low temperature of 150 °C tends the creep resistance of AZ91 to be higher than that of AE42, but at higher temperatures shows AE42 a higher creep resistance (Figure 3). That could be related to the thermal stability of the structure of AE42, otherwise the alloy AZ91 has a significant fraction of the second phase Mg17(A1,Zn)12 wich has a low melting point of 437 °C [6]. The composite material AE42-23%C has a higher improved creep resistance over the whole test temperatures than that of AZ91-23%C in case of loading parallel to the reinforced plane. 10-5

150°C

10-5 200°C

10-6

10-6

10-7

10-7

10

-8

n=13

10-9 20 dε/dtmin , s-1

8.8

10-5 10-6 10-7 10-8 10-9 20

50

12

100

21

21

16 19

50

18.5

50

100

200

AE42-C / parallel AZ91-C / parallel AE42-C / normal AE42 AZ91

n=6

7

8.6

Stress , MPa

250°C

9.8

9

10-8

10-9 200 20

12

n=9.4

100

200

Stress , MPa Figure 3: Minimum creep rate against creep stress of the Mg-alloys AE42 and AZ91 and their composite materials at 150, 200, and 250 °C.

The apparent stress exponent n of the matrix alloys was almost lower than that of the composite materials. Nearly equal high stress exponent values were attained also from creep testing of different magnesium alloys [4, 7] and magnesium composites [5, 4]. But such high stress exponent n up to 21 (Figure 3) does not provide a clear evidence about the dominant creep deformation mechanism in matrix and composite materials. So that the true stress exponent should be determined. To determine the true value of the creep stress

387 exponent n, the effective creep stress (σ-σο) should be represented against the minimum creep rate, as shown in equation [8]: ε&min =

ADGb ⎛ σ − σ e ⎞ KT ⎜⎝ G ⎟⎠

dε/dtmin , s

-1

Where σο is the threshold stress, D is the diffusion coefficient, G is the shear modulus, T is the absolute temperature, b is the burgers vector, k is the Boltzmann’s constant and A is a constant. Sklenicka et al [4] and Yong et al [8] have estimated the threshold stress (σο) for metal matrix composites by extrapolating the line fitting the minimum creep rate and the applied creep stress (Figure 3) to identify the creep stress corresponding to a mini-10 -1 mum creep rate of 10 s . -5

10

-6

10

150°C

-7

10

n≈5

-8

10

-9

dε/dtmin , s

-1

10

25

-5

10

50

75 100125

-6

n≈5

-8

10

-9

10

10

20

10

-6

10

-7

10

-8

10

-9

25

200°C

n≈5

50

75 100125

AE42-C / parallel AE42-C / normal AZ91-C / parallel AE42 AZ91

10 10

-5

Effective stress , MPa

250°C

-7

10

50

100

Effective stress , MPa Figure 4: Minimum creep rate against the effective creep stress for the matrix alloys and the composite materials.

Replotting the minimum creep rate against the effective creep stress as shown in Figure 4, stress exponent close to 5 was obtained for both the matrix and the composite materials. The true stress exponent n∼5 and the high determined threshold stress (σο = 10 to 64 MPa for the matrix and 25 to 143 MPa for the composites) propose the possible creep deformation mechanism as interaction between the mobile dislocations and the precipitates [4] occurring in the matrix alloys and the composite materials. The formation of the precipitates in the composite material AE42 was investigated using the SEM, Figure 5. It was noticed more precipitates in the composite material particularly at the interface regions in addition to that observed in the matrix. More precipitates were noticed in the crept specimen (the bright zones in Figure 5.b). EDX -analysis of the composite AE42-23%C has shown that the precipitates around the fibers have nearly the same amount of RE elements as in the precipitates away from fibers. The long time under stress and temperature has made Mn easy to diffuse and concentrate at the precipitates in the

388 crept specimen opposing to the case of the specimen tested by quasi-static stress at room temperature. Silicon and oxygen were detected in the precipitates occurring on the fibersmatrix interface, which was related to the SiO2-binder of the fibers perform leading to formation of Mg2Si and MgO which is thermodynamically favorite and is a necessary consequence of achieving wettability between the matrix and fiber [9,10].

(a)

(b)

Figure 5: SEM Micrographs for the composite AE42-23%C, after testing under (a) quasi-static stress at RT, and (b) compressive creep under 170 MPa at 150 °C for 635 h. (The bright zones represent the precipitates)

Larson-Miller-Parameter (PLM = T(C+log(tB)) representation summarizes the experimental creep results of the matrix alloys AE42 and AZ91 (Figure 6.a) and the composite materials (Figure 6.b), where T is temperature in Kelvin and tB is the life to fracture in hours, the constant C is ≈ 20 for many technical materials [11]. The creep resistance of the alloy AZ91 was higher than AE42 only at 150 °C, rather than this temperature it will be lower (Figure 6.a). The composite material AE42 loaded parallel to the reinforcing plane has the highest creep resistance over all the tested temperatures Figure 6.b. 250

250

Creep Stress , MPa

150°C

AE42-23%C / parallel

200°C 250°C

200

AZ91-23%C / parallel

200 AZ91

150 100

150 AE42

100 AE42-23%C / normal

50

50 (a)

0 8000

9000 10000 11000 12000

0 8000

(b)

9000 10000 11000 12000

PLM= T (C+log(tB))

Figure 6: Larson-Miller Parameter against the creep stress of the matrix alloys AE42 and AZ91 and their composites at 150 °C, 200 °C and 250 °C.

389

4

Conclusions

Based on the results of the creep testing of the Mg-based composites AE42-23 %C short fibers loaded in both normal and parallel to the reinforced plane, and AZ91-23 %C short fibers by loading parallel to the reinforced plane, it could be concluded that: 1. The AE42-based composite is higher creep resistant than AZ91-23 %C over the whole test temperatures (150 °C, 200 °C and 250 °C). 2. Anisotropy in the creep behavior due to distribution of the short fibers was clearly noticed by loading parallel and normal to the reinforced plane of AE42-23 %C. 3. True stress exponent n ∼ 5 for both matrix alloys and composite materials and the development of the precipitation gives a light about the possible creep mechanism that could be interaction between the mobile dislocations and the precipitates.

5

References

[1] K. U. Kainer: Partikel, Fasern und Kurzfasern zur Verstärkung von metallischen Werkstoffe, in metallische Verbundwerkstoffe, K.U. Kaíner (Herausgeber), DGM Informationsgesellschaft GmbH, Oberursel, 1994, p. 43-64. [2] C. Fritze, H. Berek, K. U. Kainer, S. Mielke, B. Weilage: Fiber Reinforced Magnesium for Automotive Applications, Magnesium alloys and their Applications, Confference, Wolfsburg, Germany, 1998, p. 635-640. [3] G. A. Chadwick, A. Bloyce: Sqeeze Cast Magnesium Alloys and Magnesium Based Composites, Magnesium alloys and their Applications, Confference, GarmischPartenkirchen, Germany, 1992, p. 93-100. [4] V. Sklenicka, M. Pahutova and T.G. Langdon: Creep of Reinfrced and unreinforced AZ91 Magnesium Alloy, Key Eng. Materials, 2000, 171-174, p. 593-600. [5] B. Sommer, K. U. Kainer, H. Berek: Kriecheigenschaften der Kohlenstoffaserverstärkten Magnesiumlegerungen AS41 und AZ91, Tagung Verbundverkstoffe und Werkstoffverbunde, DGM Informationsgeselschaft GmbH, Frankfurt, 1997, p. 513-518. [6] G. Pettersen, H. Westengen, R. Hoier, O. Lohne: Microstructure of a Pressure Die Cast Magnesium 4 wt% Aluminum Alloy Modified with Rare Earth Additions, Materials Science and Engineering, 1996, 207A, p. 115-120. [7] F. von Buch and B.L. Mordike: Microstructure, Mechanical Properties and Creep Resistance of Binary and more Complex Magnesium Scandium Alloys, Mg Alloys & their App. Conf. 98, Wolfsburg, Germany, 1998, p. 145-150. [8] Yong Li, and Terence G. Langdon: A Simple Procedure for Estimating Threshold Stresses in the Creep of Metal Matrix Composites, Scripta Materialia, 1997, 36 (12), p. 1457-1460. [9] B. Sommer, K.U. Kainer, H. Berek: Influence of the Interface on the Creep Properties of Carbon Fiber Reinforced Magnesium Alloy AS41, 11th Int. Conf. on Composite Materials, Gold Cost, Australia, 14-18 July, 1997, 3, p. 221-229. [10] Suraj P. Rawal: Interface Structure in Graphite-fiber-reinforced metal-matrix Composites, Surface and Interface Analysis, 2001, 31, p. 692-700. [11] A. Kranz: Mechanisches Hochtemperaturverhalten der Aluminiumlegierung AA2024 in Abhängigkeit von Umformprozess und Wärmebehandlungszustand, Ph.D. Thesis, LFW, RWTH-Aachen, 2002, p. 95-97.

390

High Performance Magnesium Composite Alloys with Mg2Si Dispersoids via Solid-state Synthesis Process Katsuyoshi Kondoh*, Tatsuhiko Aizawa** * Research Center for Advanced Science and Technology (RCAST), The University of Tokyo, Tokyo ** Center for Collaborative Research (CCR), The University of Tokyo, Tokyo

1

Introduction

From a viewpoint in utilizing the light weight effect of magnesium alloys to reduce both the discharge of air pollutions and the energy consumption, they should be applied to the sliding or moving automotive components such as pistons, valve lifters, valve-spools etc. In general, the magnesium alloys, however, have poor mechanical and physical properties, in particular corrosion and wear resistance, compared to the other industrial metals, e.g. steels or aluminum alloys. The previous works1-4 indicated that Mg2Si bulky intermetallic, which was synthesized in solid-state from the elemental magnesium and silicon powder mixture, had a low density of 1.88 g/cm3 and good mechanical properties, such as 350~700 Hv micro-hardness, high elastic modulus of 120 GPa and a 1085 ͠ high melting point. The salt spray test and electrochemical test (polarization curve) showed that the Mg2Si bulky material had a superior corrosion resistance to the conventional stainless steels. Accordingly, Mg2Si will conquer some defects of the conventional Mg alloys. In other words, there is a possibility to improve the poor properties by the combination of Mg alloys and Mg2Si. The useful material design is to disperse Mg2Si particles uniformly into the Mg alloy matrix (Metal Matrix Composite; MMC). In this study, the material and processing designs for Mg composites with Mg2Si dispersions are introduced, by employing in-situ formation of Mg2Si in solid-state. The effect of the refinement microstructure control on the mechanical properties of the magnesium composite dispersed with Mg2Si is discussed. Furthermore, from a viewpoint in establishing the scale-up oriented technology, the possibility to apply the processing design for this magnesium composite to the mass production is investigated.

2

Process Design for Magnesium Composite via Solid-state Synthesis

Figure 1 illustrates a schematic to fabricate the magnesium composite with Mg2Si dispersoids synthesized in solid-state from the elemental magnesium and silicon powder mixture5.

391 Input raw powder

Green billet

Pre-heating

e.g. hot extrusion

Mg alloy coarse powder

Si/SiO2 fine powder

Hot plastic workings

Mg-Si compact via RPW process

Solid-state synthesis 2Mg+Si ψ Mg2Si

Repeated Plastic Working (RPW) Feeding input materials

Compaction

Upper punchΤ Τ

Upper punchΣ Input raw materials Die Lower punch

Output green compact

Withdrawal

Backward Extrusion

Figure 1: Schematic illustration of magnesium composite alloy with Mg2Si dispersoids via repeated plastic working and solid-state synthesis.

Mg2Si intermetallic can be formed by the reaction of magnesium and silicon with an exothermic heat. Silica (SiO2) can be also supplied as another raw material to synthesize Mg2Si and MgO via the deoxidization by magnesium. Both solid-state synthesizes occur based on the following equations. 2Mg + Si → Mg2Si 4Mg + SiO2 → Mg2Si + 2MgO

(1) (2)

The Mg alloy chips or coarse powder, having the same chemical composition as the conventional one, e.g. AZ31, AM60 and AZ91, are employed as the main raw materials. The repeated plastic working (RPW) process consists of the alternative two upper punches motion for the compaction and backward extrusion as shown in Figure 16. The punch speed is about 0.5~1 m/s, and the energy into the raw material is controlled by the speed. The relative density of the green compact in the die is about 85~90%. Not only the uniform distribution of silicon particles into the magnesium powder, but also the refinement of the magnesium matrix grains and silicon particles occur by the backward extrusion. In particular, the impact energy is significantly effective on the fragmentation of brittle silicon particles. The pre-heating controls the progress of Mg2Si synthesis by using its exothermic heat. The heating conditions, in particular the temperature and holding time, are decided by the thermal analysis on the green compact after RPW, using a differential scanning calorimeter (DSC). For example, when employing the elemental AZ31-8 mass% SiO2 powder mixture, the pre-heating temperature is the ignition one (Ti) of the exothermic peak of the green compact after the RPW process as shown in Figure 2. The XRD

392 patterns in Figure 3 indicate that the magnesium completely reacted with silicon to form Mg2Si by pre-heating the green compact. Hot plastic working is necessary to consolidate the green compact into a full density. The RPWed green compact is immediately hot extruded after pre-heating to produce rods, pipes or plates, by accompanying the extrusion ratio of 37. Mg

4000

400

(b)

3500 3000

200 100

Intensity / a.u.

DTA / uV

300

Ti

0

Mg2Si

MgO

2500 2000

(a)

1500 1000

-100

SiO2

500

-200 300 400 500 600 700 800 900 1000 1100

0 20

25

Figure 2: DSC curve of RPWed green compact of AZ31-8 mass% SiO2 mixture.

30

35

40

45

50

Diffraction Angle, 2θ /°

Temperature, T / K

Figure 3: XRD patterns of input raw powder (a) and hot extruded composite (b).

3

Characteristics of Mg Composite with In-situ Synthesized Mg2Si

3.1

Microstructure control by RPW

Figure 4 shows optical microstructures with chemical etching on a transverse section of hot extruded AZ31 alloys via RPW (a) and (b), and that of AZ31 chips used as raw materials (c). After 50 cycles of RPW, the grain size of the hot extruded AZ31 alloys is 10~20μm, m, which is extremely small compared to that in employing cold compaction without such cyclical plastic workings. The grain becomes fine with the increase in the number of cycles in the RPW process. Raw chips show Mg-Al intermetallic compounds distributed in the matrix. However, they are not obviously observed in the cyclical plastic worked AZ31 alloys because of their fine dispersoids. (b) N=200

(a) N=1

(c) Raw material (AZ31)

Mg-Al intermetallic 10Ǵ Ǵm

10Ǵ Ǵm

Figure 4: Optical microstructures of hot extruded AZ31 alloys via RPW (a, b) and raw AZ31 chips.

100Ǵ Ǵm

393 Figure 5 shows microstructures of the hot extruded composites via PRW in employing AZ31-5 mass% Si powder mixture. The increase of the cycle number causes the refinement of Mg2Si dispersoids and their uniform distribution in the matrix. The AZ31 composite alloys via the cyclical plastic working process reveal a minimum particle size of less than 5 μm. (a) N=1

(b) N=50

(c) N=200

Extruding Direction

Mg2Si

10Ǵ Ǵm

10Ǵ Ǵm

10Ǵ Ǵm

Figure 5: Microstructure dependence on cycle number in RPW of hot extruded AZ31-5%Si composite.

3.2

Mechanical properties of AZ31 composites

400 350 300 250 Cold compacting (without RPW)

200 250

350

450

550

650

Matrix Grain Size, d -0.5 / m- 0.5

Figure 6: Relationship of matrix grain size and UTS of hot extruded AZ31 alloys via RPW.

Ultimate Tensile Strength, ǻ/ MPa

Ultimate Tensile Strength, ǻ / MPa

Grain refinement is one of the effectively strengthening factors in using cyclical plastic working. When employing AZ31 powder with no silicon particle as raw materials, as shown in Figure 6, the hot extruded AZ31 alloys via the cyclical plastic working process recognizes the Hall-Petch relationship between UTS (σ) and the matrix grain size (d-0.5). It indicates that the effective strengthening mechanism of the magnesium alloys via RPW strongly depends on the grain refinement of the AZ31 matrix, because the grain boundaries act as a barrier to obstruct slips of dislocations. The tensile strength of the magnesium composite strongly depends on the RPW condition as shown in Figure 7.

400

Compositions; AZ31-5mass%Si

350 300 250 200 150 100

0

50

100

150

200

Number of Cycles in RPW, N

Figure 7: Dependence of UTS on RPW cycle number in employing AZ31-5%Si.

The increase of the cycle number in RPW process causes the significant strengthening of the composite material. For example, UTS of the hot extruded composite via RPW with N = 200 cycles is 345 MPa, which is about twice as the conventional cast AZ31 alloy, when employing AZ31-5 mass% Si powder mixture. The matrix hardness also increases

394 up to 100 Hv by RPW process in increasing the cycle number. Considering the effects of microstructure refinement as mentioned above, the strengthening mechanism of the in-situ synthesized Mg2Si/Mg composite material consists of both the uniform distribution of fine Mg2Si particles and the refinement of matrix grain by RPW process.

4

Scale-up Oriented Process for Mg2Si/Mg Composites

Figure 8 shows the appearance of the 3000kN screw-driven press machine, installed in the industry, used for RPW process (a) and the AM60 green compact including 2 mass% SiO2 particles after RPW with 200 cycles by using this press (b). The diameter of the compact is 79 mm and its weight is 500 g. After pre-heating it at 753 K in nitrogen gas atmosphere, the hot extrusion is applied to fabricate the pipe and rods (showing in Figure 9) by using 5 the 8000 kN extrusion equipment . (a)

(b)

Figure 8: 3000kN press machine used for industrial RPW process (a) and green compact via RPW with 200 cycles (b) in employing AM60-2 m% SiO2 powder.

(a) Pipe

(b) Rods

Figure 9: Appearances of specimens of pipe (a) and rods (b) of magnesium composites with Mg2Si by using RPW and hot extrusion processes in manufacturing plants.

As shown in Figure 10, the same relationship between the mechanical properties of the hot extruded composites and the cycle number of RPW process is recognized, when fabricating them via scale-up oriented processes. A 5000 kN press machine indicates a signifi-

395 cant possibility to produce a large billet of about 3 kg, used for commercial products, via the RPW process for fine microstructure controls.

15

300 250

10

200

5

150 1

50

100

200

Number of Cycles in RPW㧘 㧘N

0

400

25

(b) T=753K

20

350

15

300 250

10

200

5

150 100

1

50

100

200

Elongation, ǭ / %

20

350

100

450

25

(a) T=673K

UTS㧘 㧘ǻ / MPa

UTS㧘 㧘ǻ / MPa

400

Elongation, ǭ / %

450

0

Number of Cycles in RPW㧘 㧘N

Figure 10: Dependence of mechanical properties of large-scaled magnesium composite specimens on cycle number in RPW in employing AM60-2%SiO2 powder mixture.

5

Conclusions

The characteristics of the conventional magnesium alloy are remarkably improved by the microstructure control, e.g. the uniform distribution of in-situ synthesized Mg2Si fine particles, and the refinement of the magnesium matrix grains and the dispersoids by the RPW process.

Acknowledgement This study was financially supported by the matching foundation project to create university based businesses from the METI and by the project “Environmentally Benign Processing of High-Performance Magnesium Alloys” from Kanagawa Academy Science and Technology.

References [1] G.Frommeyer, Z. Metallkd, 1994, 85, 373-377. [2] L.F.Mondolfo, Aluminum Alloys-Structures and Properties, London-Boston, Butterworths, 1976, p.566. [3] K.Kondoh, H.Oginuma, T.Aizawa, E.Yuasa, Materials Transactions, 2001, 42, 12931300. [4] K.Kondoh, H.Oginuma, A.Kimura, S.Matsukawa, T.Aizawa, Materials Transactions, 2003, 45, 981-985. [5] K.Kondoh, H.Oginuma, R.Tsuzuki, T.Aizawa, Materials Transactions, 2003, 44, 611-618. [6] K.Kondoh, T.Aizawa, Materials Transactions, 2003, 44, (in printing).

396

Fracture Toughness Behaviour of As-Cast and Long-term Annealed Pressure Die Cast Magnesium Alloys Doris Regener, Gabriele Dietze Institute of Materials Engineering and Materials Testing, Otto-von Guericke-University Magdeburg, Germany

1

Introduction

The application of pressure die cast magnesium alloys for safety relevant automotive components requires reliable values, which assess the damage tolerance of the used material. The necessity results from the fact that the manufacturing of sound die part can not be ensured. On the contrary, numerous defects like gas pores and microshrinkage endagering the component integrity are present. Futhermore, a crack initiation due to mechanical and/or thermal loading is to take in consideration, because the crack resistance can be changed due to materials deterioration. Present, there is a lack of frecture toughness data, especially for pressure die cast magnesium alloys. Fracture toughness values for cast and wrought alloys, determined on the basis of Linear Elastic Fraktur Mechanics (LEFM), are summarized in ASM Handbook 1/2 Volume 19 [1]. Typical KIc values for wrought alloys are 20–30 MPa m , while typical 1/2 values for sand cast alloys are 10–20 MPa m . Most recently, Padfield et al. [2] carried out an Elastic Plastic Fracture Mechanics (EPFM) analysis of the pressure die cast alloys AM60B, AM50A and AZ91D using the compliance unloading technique to record crack extension for earch specimen. The J–R curves generated for the AM-alloys showed stable crack propagation beyond ASTM E 1820 limits. The alloy AZ91D had unstable crack growth with a Jc of 7.5 N/mm. The purpose of this research was to determine the fracture toughness behaviour of the pressure die cast alloys AZ91, AM50 and AE42 in an as-cast and long-term annealed condition, in order to prove the influence of the microstructural changes.

2

Experimental

For the investigation plate-like components from the alloys AZ91, AM50 and AE42 (Table 1) with the dimensions of 200 x 75 x 10 mm were cast on a cold chamber die-casting machine. After that, some plates were subjected to a long-term annealing at 150 °C and 200 °C for 1000 hours. Specimens for tensile tests an fracture mechanics investigations were removed from the as-cast plates as well as from the annealed plates. Table 1. Chemical composition of alloys investigated

Alloy

%Al

%Mn

%Zn

%Si

AZ91 AM50 AE42

9.3 4.1 4.2

0.12 0.24 0.18

0.79 0.02 <0.0001 0.006 0.013 0.002

%Cu

%Ni

0.0007 0.0006 <0.0002 0.0004 0.0120 0.0009

%Fe

%RE

0.0046 0.0020 0.0023

2.614

397 For the tensile tests flat specimens with the gauge length of 30 mm were used. The fracture toughness investigations were carried out on fatigue precracked SENB(10 x 10 x 55 mm) and CT-specimens (W=20 mm, H=24 mm and B=10 mm). According to the different materials behaviour the evalutation of the results occurred on the basis of the LEFM or EPFM-analysis using for the latter the multispecimen test method. The crack extension Δa was measured on the fracture surface after heat tinting at 350 °C/2 h (AZ91) and 430 °C/2 h (AM50 and AE42). For the investigation of the microstructure by means of light an scanning microscopy (SEM) the preparation of specimens consisted of wet grinding until 4000 mesh silicon carbide paper, polishing with finally 1 μm diamond paste and etching with 3% alcoholic nitric acid. The fracture appearances were evaluated on the basis of SEM fractographs.

3

Results and discussion

3.1

Microstructure

The metallographic investigation of the as-cast alloy AZ91 reveals a cored structure with primary α-Mg solid solution (1, 2) surrounded by a divorced eutectic of α-Mg and βphase Mg17Al12 (3) (Fig. 1a). EDS-analysis shows that the eutectic α-Mg is Al richer (9– 11 wt. %) than the primary α-Mg(4–6 wt %), what is also visible due to the different etching characteristics of th phases. During the long-term annealing at 150°C further β-phase precipitates mainly from the Al rich eutectic α-Mg in an discontinuous kind, which crystallizes to the already existing massive β-phase or forms isolated round or lamellar particles (Fig. 1b). The 1000 h – annealing at 200 °C shows the same tendency of the precipitation. However, because of the higher temperature a distinct coagulation and coarsening (1) of the precipitates were observed (Fig. 3c). Dargusch et al. [3, 4] found such extensive precipitation of ß-Mg17 Al12 after creep tests of this alloy at temperatures between 125 and 175 °C. The fact that the long-term annealing brings out the same microstructural changes reveals the great role of the high-temperature exposure.

Figure 1. SEM micrographs of the alloy AZ91 in the as-cast condition (a) and after 1000 h long-term annealing at 150 °C (b) and 200 °C (c)

In comparison to AZ91. the as-cast microstructure og the alloy AM50 precipitates les β-phase due to the reduced Al-content. In addition, because of the higher Mn-content some precipitates of the type Al8Mn5 are found. The precipitation process due to the longterm annealing is also comparable with that of the alloy AZ91. However, the precipitation

398 is not so extensively and the precipitation region contains in addition to β-phase some Al8Mn5-particles. The as-cast microstructure of the alloy AE42 consists of cored globular, partly long α-Mg grains and eutectic precipitates in lamellar, rarely in globular shape. These precipitates, which are of Al-Re type corresponding to an EDS-analysis, are arranged at the boundaries of the α-Mg grains. During the long-term annealing at 150 and 200 °C the microstructure behaves rather stable. Only the precipitation of less globular or plate-like particles, which are rich with Al and the rare earths Ce, Nd an Pr, can be proved. 3.2

Tensile tests

The mechanical properties of a die casting depend strongly on the fabrication route and fabrication variables involved. The effect of these variables, which determine the size and distribution of gas pores and microshrinkage, might be greater than the effect of alloying element concentration. Therefore, the tensile properties required for the evaluation of the fracture toughness investigation were measured on specimens of the same dimensions taken from the plates. The average values of the yield strength Rp0.2, ultimate strength Rm and strain to fracture A30 mm are given in Table 2. Table 2. Tensile properties at room temerature of speximens taken from plates

Alloy AZ91

AM50

AE42

Condition as-cast 150°C/ 1000 h 200°C/ 1000 h as-cast 150°C/ 1000 h 200°C/ 1000 h as-cast 150°C/ 1000 h 200°C/ 1000 h

Rp0.2 [N/mm2] 118 163 105 83 123 115 92 127 125

Rm [N/mm2] 180 183 151 163 181 157 174 170 154

A30 mm [%] 2.1 0.4 1.3 4.5 3.4 2.7 4.2 2.9 2.0

The distinctly less strength and stain values of the as-cast condition in comparison to separatly cast speciemens [5, 6] are essentially attributed to the higher number of internal manufacture-induced defects in the thicker plates. The long-term annealing at 150 °C/100 h leads to a considerable increase of the Rp0.2-value due to the numerous, more or less fine precipitates in the eutectic α-phase. The alloy AZ91 shows the highest increase because of the high precipitation density of Mg17Al12-particles. Although the microstructural change in the alloy AE42 was less, as proved by means of SEM, an increase of the yield strength occurred, too. With regard to the ultimate tensile strength, only the alloy AM50 shows an increase due to the long-term annealing at 150 °C. It is assumed that because of the porosity the specimens do not always reflect the true intrinsic materials properties. The decrease of the Rp0.2-values during the long-annealing at 200 °C, especially of the alloy AZ91, is connected with the coagulation and coarsening ofprecipitates. From the reduced Rm-values results that due to the suoerimposed effect of microstructural degradation and manufacture-induced defects the materials is not able to absorb high stresses. The strain to fracture is already low in the as-cast condition, what shows the high sensitivity of the material to cast-induced porosity. Due to the long-term annealing it decreases yet more provided that the precipitates block effectively the plastic deformation.

399 3.3

Fracture toughness tests

The described changes of the microstructure due to the long-term annealing results in a different fracture toughness behavior. The strongest changes occur in the alloy AZ91 and lead to an unstable crack growth of the specimens. Therefore, with regard to all AM50and AE42-specimens and the as-cast condition of AZ91 the multiple specimen method is used to measure stable crack growth and to plot crack resistance curves J-Δa, while the parameters of the LEFM are required for the evaluation of the annealed conditions of AZ91. Selected load-displacement-curves obtained from long-term annealed CT-specimens of the alloy AZ91 and AE42 show already distinct differences of the energy absorption (Fig. 2). Because of its high embrittlement the alloy AZ91 behaves essentially elastic and the energy absorption capability is strongly limited. Exceeding the load maximum the crack grows unstable. In opposite, the deformation behaviour of the alloy AE42 is elasticplastic with a high plastic part. The fatigue precrack can progress in a stable kind.

Figure 2. Force-displacement curves obtained from CT-specimens

A comparison of the crack resistance behaviour of all three material conditions on the basis of J-Δa-plots was only possible or the alloys AM50 and AE42 using CT-specimens. The curve fit was done with a power function of the type J=A ⋅ ΔaD, in which the exclusion of Δa-values beyond Δamax was considered [7]. The exclution line of Jmax was not reached. In order to characterize the crack initiation, the crack resistance J0.2 is used. The dermination of Ji-values is not possible, because measurable stretch zones can not be found. The evaluation of the resistance to crack growth occurs on the basis of the slope of curve ΔJ/ Δa for a comparable Δa-region of 0.4 mm within the valid curve limit. Respecting the alloy Am50 (Fig. 3) the J-values of the ass-cast condition exceed considerably that of the annealing states. The J0.2-value of this as-cast condition amounts 20.1 N/mm and is the highest of all (Table 3), what also reflects the usually well-known ductile behavior of Al-Mn-alloys. The resistance to crack growth is most impaired by manufactureinduced internal defects. The deterioration of the crack initiation behaviour after the lonterm annealing is attributed to the increased precipitation density of the brittle intermetallic phase Mg17Al12. Because of the scattering of the measured J-Δa-pairs distinct differences between the both annealed conditions are not recognizable. The higher slope of curve characterizing the annealing at 200 °C is explained by the coagulation of the β-phase.

400

Figure 3. Crack resistance curves of the alloy AM50 obtained from CT-specimens

In contrast to AM50, the alloy AE42 shows the lowest values of crack initiation and growth in the as-cast condition (Fig. 4). Due to the long-term annealing an improvement of the values occurs. This is attributed to the homogenization of the material without essential microstructural changes, which point to the high thermal resistance of AE42 in comparison to the alloys AZ91 and AM50.

Figure 4. Crack resistance curves of the alloy AE42 obtained from CT-specimens

Respecting the alloy AZ91 the strong changes of the microstructure due to the longterm annealing result in the distinct decrease of the fracture toughness. While for the as-cast specimens a stable crack growth occurred and a J-Δa curve could be describe, the annealed specimens failed unstable and the evaluation was carried out according to the LEFM-recommendation [8]. However, because of the limited dimensions of the used CTand SENB-specimens the most measured values are only to define as KQ, not as valid KIc (Table 3). The SEM fractographs of CT-specimens selected for the long-term state 150 °C/1000 h show the different deformation behavior of the alloys. The fracture surface of AZ91 assumes the features of the cleavage fracture characterized by cleavage facets, steps and secondary cracks (Fig. 5a). This fracture mode is connected with the failure of the α-Mg/Mg17Al12 interface and the cleavage of Mg17Al12-particles. The fracture surfaces of

401 the AM50- and AE42 specimens (Fig. 5b and c) appear rougher and a tearing of the material is visible. The middle area of the fractograph of AE42 (Fig. 5c) shows probably the deformation features due to twinning. Table 3. Summary of the fracture toughness values

Alloy Condition

AZ91 as-cast 150 °C/1000 h 200 °C/1000 h AM50as-cast 150 °C/1000 h 200 °C/1000 h AE42 as-cast 150 °C/1000 h 200 °C/1000 h

J0.2[N/mm]

ΔJ/Δa

CT 12.1

SENB 4.8

CT 8.3

20.1 12.4 11.5 13.5 14.2 15.8

10.2

12.3 14.6 19.6 10.3 21.7 26.6

7.4 12.7 11.5

KIc KQ [MPam1/2] [MPa m1/2] SENB CT SENB CT 4.0 9.9-11.8 7.4-10.4 9.9 8.9-11.8 9.5-10.8 4.6 2.7 22.2 17.4

Figure 5. SEM fractographs of CT-specimens from AZ91 (a), AM50 (b) and AE42 (c) in the long-term annealed condition 150 °C/1000 h

4

References

[1] ASM Handbook Volume 19: Fatigue and Fracture, ASM International, Materials Park, OH, 877–878. [2] C. J. Padfield, T. V. Padfield in Light Metals for the Automotive Industry, SAE SP1683, Warrendale, PA, USA, 2002, 27–33. [3]. M. Dargusch, M. Hisa, C. H. Cáceres, G. L. Dunlop in Proceedings of the Third Int. Magnesium Conference (Ed.: G. W. Lorimer), Inst. of Metals, London, UK, 1996, 153–165. [4} M. S. Dargusch, G. L. Dunlop in Magnesium Alloya and their Applications (Eds.: B. L. Mordike, K. U. Kainer) Werkstoff-Informationsgesellschaft mbH, Frankfurt, Germany, 1998, 277–282. [5] Hydro Magnesium: Datenblatt Magnesium-Druckgusslegierungen [6] D. Regener, G. Dietze, H. Heyse, Mat.-wiss. u. Werkstofftech. 2002, 33, 606–613. [7] ISO/FDIS 12135, Final draft 2002 [8] DIN EN ISO 12737, April 1999

402

Influence of Strain Rate and Temperature on Deformation and Fracture Behaviour of Magnesium Alloy MgAl8Zn: Tests and Numerical Simulations E. El-Magd and M. Abouridouane Department of Materials Science, RWTH Aachen

Abstract The flow and fracture behaviour of the wrought magnesium alloy MgAl8Zn was investigated by quasi-static and dynamic uniaxial compression and tension tests in the strain rate range of 0.001 s −1 ≤ ε& ≤ 5000 s −1 and temperatures between 20 °C and 300 °C. Shear tests with hat shaped specimens were carried out by quasi-static and dynamic loading in the shear rate range of 0.01 s −1 ≤ γ& ≤ 116000 s −1 and a temperature of 20 °C. Using of experimentally determined stress-strain curves, the effect of strain, strain rate and temperature (strain hardening, strain rate sensitivity and thermal softening) on the stress and the compressive deformation at fracture was determined, showing that both flow stress and the limits of possible deformation are controlled by strain rate and temperature. The mechanical flow behaviour of the magnesium alloy MgAl8Zn under compression loading at high strain rates was characterised with a constitutive material law under assumption of the domination of damping controlled glide processes, taking the adiabatic behaviour of the deformation process into consideration. Compression, tension and shear tests were simulated by the finite element method in order to validate the used constitutive material law. Metallographic investigations on hat shaped specimens showed shear localization by dynamic loading, which is consent with the numerical simulation results.

1

Introduction

Forming and machining of brittle materials such as magnesium alloys with conventional methods is limited, so that an intermediate treatment of the work piece should be necessary. Under dynamic loading, metallic materials exhibit an increase of flow stress and ductility with increasing deformation rate [1]. Materials of lower ductility can be deformed to higher strains by using high strain rate deformation processes such as magnetic or explosive forming [2,3]. Therefore, adequate knowledge is needed in this field. Furthermore, constitutive material laws are necessary to describe the material behaviour in simulation of high strain rate deformation processes as crash test and high speed cutting. The strain rate and temperature influence considerably the material flow behaviour in deformation processes. Under quasi-static loading strain hardening causes an increase of force and acts stabilising on the deformation process. In case of dynamic loading, additional influences on the flow stress and the ductility of the material have to be taken into

403 consideration. With increasing deformation rate, the strain rate sensitivity increases leading to a higher value of flow stress and stabilises the deformation. On the other hand, the adiabatic character of the deformation process reduces the flow stress and promotes instability. Furthermore, the deformation process is influenced by inertia [4] and mechanical wave propagation effects [5]. The aim of this work is to determine the influence of strain rate and temperature on flow and fracture behaviour of wrought magnesium alloy MgAl8Zn subjected to compression, tension and shear loading.

2

Experimental Investigation

The deformation and fracture behaviour of the high-strength wrought magnesium alloy MgAl8Zn (AZ80: 8.3 % Al, 0.6 % Zn) was investigated by quasi-static and dynamic uniaxial compression and tension tests (cylindrical specimens) in the strain rate range of 0.001 s −1 ≤ ε& ≤ 5000 s −1 and temperatures between 20 °C and 300 °C. Shear tests with hat shaped specimens were carried out by quasi-static and dynamic loading in the shear rate range of 0.01 s −1 ≤ γ& ≤ 116000 s −1 and a temperature of 20 °C. The investigated specimens were machined with their load axis parallel to the extrusion direction of the test material in as-delivered condition. For strain rates ε& ≤ 100 s-1, the tests were carried out using a numerical controlled hydraulic testing machine. High strain rate experiments with ε& ≥ 1000 s-1 were performed on a Split Hopkinson Pressure Bar. Experimentally determined stress-strain curves of the investigated Mg alloy MgAl8Zn are represented in fig. 1. The stress of the quasi-static flow curves increases continuously with increasing strain due to strain hardening. At lower strains, the stress-strain curves show an increase of flow stress with the strain rates, on the other hand a decrease of stress beyond a stress maximum due to thermal softening caused by the adiabatic character of the deformation process resulting in a temperature increase in the specimen at high strains. 500

Compression Tests at 20°C MgAl8Zn

300

Nominal Stress, MPa

Flow Stress, MPa

400

Strain Rate -1

4659 s -1 3997 s -1 3300 s -1 2200 s -1 2096 s -1 10 s -1 0.01 s

200 100 0

0

0.05

0.10

0.15

0.20

200

Tension Tests at 20°C MgAl8Zn

400 300

Shear Stress, MPa

500

Strain Rate -1

2045 s -1 2030 s -1 1975 s -1 1773 s -1 100 s -1 10 s -1 0.01 s

200 100 0

0

5

10

15

20

Shear Tests at 20°C MgAl8Zn

150 Shear Strain Rate -1

116000 s -1 99000 s -1 97000 s -1 96000 s -1 88000 s -1 173.2 s -1 17.32 s -1 0.0173 s

100

50

0

0

0.5

1.0

1.5

2

Figure 1. Stress-strain curves of AZ80 in compression, tension and shear tests at different strain rate and 20 °C

Fig. 2. shows the strain rate and temperature dependency of the flow stress (at a true plastic strain of ε = 0.01) and of the compressive deformation at fracture. The flow stress of MgAl8Zn increases with the strain rate, whereby this increase becomes more significant at lower strain rates and higher temperature or at higher strain rates and lower temperature (fig. 2a.). In the investigate strain rates range, three different regions can be dis-

404 tinguished. In the region of low temperatures and moderate strain rates, only a slight dependence of the flow stress on strain rate and temperature is observed. The usual plastic deformation mechanism dominates in this region. At temperature 200 °C and higher, a greater influence of strain rate and temperature can be observed at low strain rates. This is caused by stress relaxation due to the creep deformation mechanism. At high strain rates ( ε& ≥ 1000 s −1 >), the viscous damping deformation mechanism can be considered as dominant. 35

Flow Stress, MPa

300

300 °C 200 °C 100 °C 20 °C

200

100 Creep Deformation Mechanism

(a)

0 -3 10

Plastic Deformation Mechanism Damping controlled Mechanism -1

10

1

3

10

10 -1

Strain Rate s

Deformation at Fracture εf, %

Temperature Compression Tests MgAl8Zn ε = 0.01

Compression Tests MgAl8Zn Temperature 300°C 250°C 200°C 150°C 100°C 20°C

25

15

(b) 5 -3 10

-1

10

1

3

10

10 -1

Strain Rate s

Figure 2. Influence of strain rate and temperature on the stress and the compressive deformation at fracture

The deformation at fracture ε f , defined as the relative reduction of the height of the cylindrical compression specimen at fracture is considered as a measure for the compressive ductility of MgAl8Zn. The deformation at fracture is determined as a function of strain rate for different temperatures as represented in fig. 2b. At room temperature with increasing strain rate in the range of 0.01 s −1 ≤ ε& ≤ 1000 s −1 the deformation at fracture shows an increase due to the stabilizing effect of the strain rate sensitivity for MgAl8Zn. Under impact loading with strain rate higher than 1000 s-1, the ductility of MgAl8Zn increases sharply with the strain rates due to higher strain rate sensitivity of the material, whereby this dependence were clearly influenced by temperature. Under quasi-static loading with ε& = 0.001 s −1 , the fracture strain increases with the temperature. No fracture was detected at temperatures of 200 °C and higher up to values of ε = 80 %. For the strain rates of ε& = 0.01 s −1 and ε& = 0.1 s −1 the absence of fractures was observed at temperatures of 250 °C and higher.

3

Constitutive Material Law

In case of dynamic loading with strain rates higher than 1000 s-1, the influence of strain rate can be represented by a linear relation σ = ση + ηε& according to the damping mechanism, as it was confirmed e.g. in [6]. The parameter σh signifies the stress extrapolated from the range of high strain rates down to ε& = 0 s−1 ; η is the damping parameter. The influence of a temperature increase on the flow stress can be considered by a multiplicative function [7], so that the material behaviour can be described by [8]

405 ⎡ (Τ − Τ0 ) ⎤ σ = ⎡⎣ Κ(Β + ε)n + ηε& ⎤⎦ exp ⎢ −β Τm ⎥⎦ ⎣

(1)

where B, K and n are material constants. Tm, T0 and T are the absolute melting point, room temperature and actual temperature respectively. β is a material constant which can be set to 3 for several materials [8]. Assuming that the major part of deformation energy is transferred to heat during a dynamic deformation process and that the remaining part is consumed by an internal energy increase e.g. due to dislocation multiplication. The temperature increase can be determined by ρ c dT = κ σ dε . In this equation κ = 0.9 [9] represents the fraction of energy transferred to heat, ρ and c are the density and the specific heat capacity of the material. With the relation between flow stress and temperature according to eq. (1), the increase of temperature can be determined as a function of strain by integration. The stress-strain relation under consideration of the two contrary influences of strain rate sensitivity and thermal softening results from the substitution of temperature in eq. (1): σ=

Κ(Β + ε)n + ηε&

(2)

exp [β(Τs − Τ0 )/ Τm ] + a ⎡⎣Κ(Β + ε)n + ηε& ⎤⎦ ε& dt

∫

with Ts is the test start temperature and α = κβ/( Tm rc ) . Fig. 3. shows a comparison between computational results following eq. (2) (continuous curves) and experimental results from quasi-static and dynamic tests (markers) at different strain rates and 20 °C. The parameters determined for MgAl8Zn are given in fig. 3. A procedure for its systematic determination is illustrated in reference [10].

Flow Stress, MPa

500

500

dynamic Compression Tests ϑ = 20°C

400

400

300

0

0.04

0.08

0.12

300

200

100 0.16 0

quasi-static Tension Test -1 dε/dt = 0.01 s and ϑ = 20°C

400

300 Material Law: Strain Rate -1 K = 750 MPa 4600 s -1 B=0 4100 s -1 n = 0.3 3000 s -1 η = 0.025 MPa s 2100 s -1 a = 0.00166 MPa

200

100

500 Experiment Material Law K = 670 MPa B = 0.0047 n = 0.27

Experiment Material Law K = 605 MPa B = 0.008 n = 0.192

200 quasi-static Compression Test -1 dε/dt = 0.01 s and ϑ = 20°C 0.02

0.04

0.06

0.08

100

0

0.05

0.10

0.15

0.20

Plastic Strain

Figure 3. Constitutive description of dynamic and quasi-static flow curves of MgAl8Zn

4

Numerical Simulation of Compression, Tension and Shear tests

Compression, tension and shear tests on MgAl8Zn at room temperature were simulated by the finite element method. The material behaviour in the quasi-static (ABAQUS\Standard, axisymmetrical) and the dynamic (ABAQUS\Explicit, axisymmetrical) simulations was described by the constitutive material law eq. (2). In both cases, the numerical simulation was calculated up to the experimentally determined global deformation at fracture consid-

406 ering Coulomb friction with a coefficient μ = 0.1. Fig. 4. shows the idealization of the investigated specimens with boundary conditions. The motion of the bars was given by measured data from the experiment. In case of compression and shear tests, the output bar was idealized with infinite elements, so no reflection of mechanical waves occurs during the simulation. Compression test

Tension test

Shear test

load contact

friction μ=0.1

contact

contact symmetry axis output bar infinite elements

infinite elements

Figure 4. Idealization of the investigated compression, tension and shear specimens

A comparison between the numerical simulation results (continuous curves) and experiment (markers) is represented in fig. 5. The deviation of the computation curves from the experimental data at high strains indicates successive damage in connection with the starting fracture process, which was not considered in the material characterisation. 500

300

400 3000 s -1

300

0.01 s

Experiment Simulation Experiment Simulation

200

100

-1

0

0.05

0.10

Plastic Strain

0.15

Nominal Stress, MPa

Flow Stress, MPa

Compression Tests ϑ = 20°C

Shear Tests ϑ = 20°C

quasi-static Tension Test -1 dε/dt = 0.01 s and ϑ = 20°C

400

Shear Stress, MPa

500

300

Experiment Simulation

200

100

0

5

10

15

Plastic Nominal Strain, %

20

200

116000 s

100

0

0.0173 s

-1

-1

Experiment Simulation Experiment Simulation 0

0.5

1.0

1.5

2.0

Plastic Shear Strain

Figure 5. Comparison of experiment and numerical FE-simulation of compression, tension and shear tests

Fig. 6. represents optical micrograph (Longitudinal sections) of hat shaped specimens deformed up to the fracture under quasi-static ( γ& = 0.0173s −1 , ϑ = 20°C ) and impact ( γ& = 116000s −1 , ϑ = 20°C ) loading as well as the results of its numerical simulations (shear stress and shear strain fields). In both cases, the investigated specimens come to ductile shear fracture, whereby a shear localization was observed by dynamic loading, fig. 6f. In quasi-static simulation, the failure zone of the specimen exhibits relatively diffuse shear strain distribution and does not appear to localize, fig. 6d. Oppositely by dynamic loading, the deformation is concentrate due to thermal softening of the material and higher

407 values of shear strain are reached (fig. 6h.), which is in agreement with the experimental results. γ& = 0.0173 s −1 , γ plastic = 1.7

γ& = 116000 s −1 , γ plastic = 2.0

(a)

(e)

(b)

(f )

(c) τ [Pa ]

(d)

γ [ −]

(h)

(g) τ [Pa ]

γ [ −]

Figure 6. Shear strain localization by impact shear test (ϑ = 20 °C)

5

Conclusions

The flow and fracture behaviour of the wrought magnesium alloy MgAl8Zn was investigated by quasi-static and dynamic uniaxial compression and tension tests in the strain rate range of 0.001s −1 ≤ ε& ≤ 5000s −1 and temperatures between 20 °C and 300 °C. Shear tests with hat shaped specimens were carried out by quasi-static and dynamic loading in the shear rate range of 0.01s −1 ≤ γ& ≤ 116000s −1 and a temperature of 20 °C. Using of experimentally determined stress-strain curves, the effect of strain, strain rate and temperature (strain hardening, strain rate sensitivity and thermal softening) on the stress and the compressive deformation at fracture was determined, showing that both flow stress and the limits of possible deformation are controlled by strain rate and temperature. The mechanical flow behaviour of the magnesium alloy MgAl8Zn under compression loading at high strain rates was characterised with a constitutive material law under assumption of the domination of damping controlled glide processes, taking the adiabatic behaviour of the deformation process into consideration. Compression, tension and shear tests were simulated by the finite element method in order to validate the used constitutive material law. Metallographic investigations on hat shaped specimens showed shear localization by dynamic loading, which is consent with the numerical simulation results.

408

6

References

[1] [2] [3] [4] [5] [6] [7] [8] [9] [10]

H. Lips, M. Brodmann, E. El-Magd, J. Physique IV 10 (2000) 371 L. W. Meyer, Tagungsband Werkstoffprüfung 2000 in Bad Nauheim, 63 A. Neubauer, H. Stroppe, H. Wolf,Verlag Technik, Berlin 1988 E. El-Magd, C. Treppmann, H. Weisshaupt, J. Phys IV, suppl. III, 1997, 7, 511 H. Lippmann, Springer Verlag, Berlin, 1981, 210 K. Sakino, J. Shiori, J. Phys. IV, suppl. III, 1991, 1, C3-35 N. J. Petch, Phil. Mag. 1958, 8, 1089 E. El-Magd, J. Phys. IV, suppl. J. Phys. III 1994, 4, 149 G. I. Taylor, H. Quinney, Proc. R. Soc. A413 (1934) 307 M. Brodmann, PhD-Thesis, RWTH Aachen, Germany, 2001.

409

Low-Temperature Superplasticity in Microcrystalline ZK60 and AZ91 Magnesium Alloys Processed by ECAP V. N. Chuvil’deev1, T. G. Nieh2, M. Yu. Gryaznov1, A. N. Sysoev1, V. I. Kopylov3 1

Research Physical & Technical Institute of Nizhny Novgorod State University, Nizhny Novgorod, Russia Lawrence Livermore National Laboratory, University of California, Livermore, USA 3 Physical - Technical Institute of National Academy of Science, Minsk, Belarus 2

1

Introduction

Magnesium alloys are one of the lightest structural materials. However, they have relatively low plasticity, which causes difficulties for a wide application of magnesium alloys. An elongation to failure is usually less than 20 % for commercial cast magnesium alloys at room temperature and is about 30 - 40 % at 300 °C [1]. A practical way to enhance ductility of alloys is to create a microstructure that would facilitate the occurrence of superplasticity. In this regard, low-temperature superplasticity of magnesium alloys (from room temperature to 300 °C) is of special interest. Low-temperature superplasticity has been demonstrated to occur in many alloys with non-equilibrium microcrystalline structures processed by equal-channel angular pressing (ECAP) [2-4]. Particularly, lowtemperature superplasticity has been reported in some microcrystalline magnesium alloys processed by ECAP at relatively low strain rates (10-4 - 10-5 s-1) [5-12]. For commercial interest, however, it is desirable to obtain superplasticity in high-strength magnesium alloys not only at low temperatures but also high strain rates (> 3 × 10-3 s-1). To achieve this goal for microcrystalline high-strength ZK60 and AZ91 magnesium alloys, two major tasks are encountered. The first task is to optimize the ECAP process to develop ultrafine grains with non-equilibrium grain boundaries [2, 4]. The second task is to characterize the superplasticity properties and to obtain maximum elongation.

2

Experimental

Two magnesium alloys: ZK60 (Mg – 6 wt.% Zn – 0.5 wt.% Zr) and AZ91 (Mg – 9 wt.% Al – 1 wt.% Zn – 0.2 wt.% Mn), both in cast and microcrystalline states, were studied in the present work. Cast alloys have microstructures with inhomogeneous grain size distribution, and the mean grain sizes in cast ZK60 and AZ91 alloys were 20 and 15 μm, respectively. Microcrystalline alloys were obtained by ECAP processing of cast samples with a size of 22 × 22 × 160 mm. Multiple pressings through a channel die with a 90° die angle were performed 6 times with a velocity of 0.4 mm/s at temperatures between 180 and 230 °C for ZK60, and between 200 and 250 °C for AZ91 using the socalled “route Bc” process [13]. ECAP produces a homogeneous microcrystalline structure with a mean grain size ~ 1 μm in ZK60 and 0.8 μm in AZ9, respectively.

410 The investigation of high strain rate superplasticity was carried out using an automated complex 2167 P-50 (analogous to Instron machine). Tensile samples were mechanically polished before testing and then spark eroded to produce a square cross-section of 2 × 2 mm and a gauge length of 3 mm. Superplasticity tests were carried out in the temperature range of 20 - 430 °C in an air furnace and initial strain rate ranging from 6 × 10-4 to 3 × 10-1 s-1. To attain thermal equilibrium, samples were held at the testing temperature for 400 s before testing.

3

Results

Nominal stress, MPa

Tensile stress-strain curves for microcrystalline ZK60 and AZ91 magnesium alloys at temperatures between 200 and 430 °C and a strain rate of 3 × 10-3 s-1 are shown in Figure 1 (a and b). An increase in ductility and decrease in flow stress is observed between 200 and 260 °C. A temperature of about 260 °C is apparently the optimum test temperature for microcrystalline ZK60 as its stress - strain curve exhibits the lowest flow stress and maximum ductility. A higher flow stress and lower ductility at 280 °C is probably resulted from the fact that excessive grain growth takes place in microcrystalline ZK60 at T > 260 °C. A conclusion can be drawn in microcrystalline AZ91 alloy at the temperature of 300 °C. ECAP ZK60, 0.003 1/s

40 30

280 оС

о

200 С

20

о

240 С

10

о

430 С

0 0

100

200

a Nominal stress, MPa

о

260 С

300

400

500

600

Nominal strain, %

60

700

800

ECAP AZ91, 0.003 1/s

50

о

200 С

40 о

250 С

30

о

300 С

20 о

350 С

10

о

400 С

0 0

100

200

300

400

Nominal strain, %

500

600

b Figure 1: Nominal stress-strain relations at strain rate of 3 × 10 s for microcrystalline ZK60 (a) and AZ91 (b) alloys. -3

-1

411

Elongation to failure,%

Figure 2 displays the elongation to failure, δ, as a function of deformation temperature for both microcrystalline ZK60 and AZ91. As shown in the figure, tensile elongation as a function of test temperature at the constant strain rate of 3 × 10-3 s-1 is non-monotonic. In the microcrystalline ZK60 alloy, the elongation increases rapidly from 170 to 810% as the deformation temperature increases from 150 to 260 °C. However, it decreases sharply down to 170% at 280 °C (curve 1 in Figure 2). A further increase in deformation temperature to 430 °C causes a slightly increased plasticity to 340%. A similar behavior is observed in the microcrystalline AZ91 alloy where δ increases monotonically from 30 to 570 % as the temperature increases from 20 to 300 °C (curve 2 in Figure 2). Above 325 °C, however, δ decreases rapidly to 350 %. 1000

ECAP ZK60 (0.003 1/s) ECAP AZ91 (0.003 1/s)

800

1

600

2

400 200 0 0

100

200

300

Temperature, oС

400

Figure 2: Elongation to failure as a function of deformation temperature for microcrystalline ZK60 (1) and -3 -1 AZ91 (2) alloys at a strain rate of 3 × 10 s .

1200 1000

ECAP ZK60 (elongation) ECAP ZK60 (UTS)

84

ECAP AZ91 (elongation) ECAP AZ91 (UTS)

70

800

56

600

42

400

28

200

14

0 0,0001

UTS, MPa

Elongation to failure,%

The elongation to failure and ultimate tensile strength as functions of strain rate ε& at the optimal superplasticity temperature (260 °C for ZK60 and 300 °C for AZ91) are presented in Figure 3. With a decreasing strain rate, the values of δ increase and reach 960% and 570 % in ZK60 and AZ91 alloys at ε& = 6.5 × 10-4 s-1 and ε& = 3×10-3 s-1, respectively.

0 0,001

0,01

Strain rate, s-1

0,1

1

Figure 3: Elongation to failure and ultimate tensile strength (UTS) as functions of strain rate for microcrystalline ZK60 (260 °C) and AZ91 (300 °C) alloys.

412

4

Analysis

A comparison between the data obtained in the present work and those from previous studies [1, 6-8] is summarized in Table 1. The data show that the ductility of microcrystalline ZK60 and AZ91 alloys at temperatures range of 200 - 300 °C and at a strain rate of -3 -1 about 10 s is almost twice higher than that for similar alloys not produced by ECAP. It is also evident in Table 1 that ECAP not only enhances high temperature plasticity in ZK60 and AZ91 but also improves the room temperature properties of the alloys. For instance, elongation to failure for the microcrystalline ZK60 alloy at room temperature averages about 45%, which is more than twice the value for the same alloy processed differently [1]. Also, the microcrystalline AZ91 alloy shows a significantly improved 30% elongation at room temperature and a strain rate 3 × 10-3 s-1 (compared to merely 5 - 10 % in cast AZ91 [10]). The improved room temperature ductility has been suggested to be a result of texture randomization [12]. It is especially noted that with the ductility improvement the strength of ECAP-processed materials at room temperature remained essentially the same. Texture randomization results in a strength decrease [12], but grain size reduction causes a strength increase via the Hall-Petch strengthening mechanism. Apparently, these two factors balance out. From an application point of view, the effect of ductility enhancement at room temperature without a significant decrease in strength (Table 1) is obviously very desirable. Table 1. Tensile properties at room temperature and temperatures between 150 and 325 °C for ZK60 and AZ91 magnesium alloys produced by different techniques.

Alloy

Production Strain method rate, 1/s

ZK60

ECAP ECAP

ZK60

Deformation Elongationtemperature, to-failure, °C %

Ultimate Ref. tensile strength, MPa

3 × 10-3

260

810

12

This work

-3

200

420

30

[7]

-3

1 × 10

ZK60

Extrusion

1 × 10

325

400

16

[6]

ZK60

Forging

-

150

36

190

[1]

-3

ZK60

ECAP

3 × 10

20

45

260

This work

ZK60

Extrusion

-

20

18

370

[1]

ZK60

Casting

-

AZ91 AZ91 AZ91 AZ91

ECAP ECAP ECAP Extrusion

20

5

275

[1]

3 × 10

-2

250

375

55

This work

3 × 10

-3

300

570

23

This work

6 × 10

-5

200

660

25

[8]

-5

200

80

100

[8]

-3

20

30

250

This work

6 × 10

AZ91

ECAP

3 × 10

AZ91

Casting

-

20

3

235

[10]

AZ91

Extrusion

-

20

13

341

[10]

413

5

Discussion

The occurrence of superplasticity in microcrystalline magnesium alloys at relatively low temperatures of 200 - 300 °C and strain rate of 3 × 10-3 s-1 can be rationalized as follows. The constitutive law for superplastic flow is given by [14] n

p

⎛ σ ⎞ ⎛ b ⎞ ⎛ GDb 0ϕ ⎞ ⎛ Q ⎞ ε& = A ⎜ ⎟ ⎜ ⎟ ⎜ ⎟ exp ⎜ − RT ⎟ ⎝ G ⎠ ⎝ d ⎠ ⎝ kT ⎠ ⎝ ⎠

(1)

where b is Burgers vector, d is mean grain size, ϕ is thickness of grain boundary, Db 0 is pre-exponential factor of grain-boundary diffusion, G is shear modulus, k is Boltzmann constant, Q is activation energy of grain-boundary diffusion, R is molar gas constant, p = 2 and A = 102 are the numerical factors. The parameter n is the reciprocal of strainrate sensitivity coefficient ( m = ∂ ln σ ∂ ln ε& ). The n values can be determined from the analysis of the experimental dependence σ ( ε& ) , and they are n = 2.5 (m = 0.4) and n = 3.3 (m = 0.3) for microcrystalline ZK60 and AZ91 alloys, respectively. From Equation (1) the activation energies can be determined to be 75 and 78 kJ/mol for microcrystalline ZK60 and AZ91 alloys, respectively. These values are much lower than the reported value of 92kJ/mol for the grain boundary diffusion in Mg with equilibrium grain boundaries [15]. For example, the grain-boundary diffusion coefficient Db* = Db 0 exp(−Q / RT ) at temperature of 260 °C is 25 - 40 times higher than the value of Db = 7.4 ⋅ 10-12 m2/s calculated on the basis of Q = 92 kJ. We have developed a model to explain enhanced grain boundary diffusion in a material with a non-equilibrium grain boundary structure during superplastic deformation [4, 16]. In accordance with the theory [16], a value of Db* ( ε& ) can be written as: Db* ( ε& ) = Db (1 + M1 ln (1 + M2ε& ) ) ,

(2)

where parameters M1 and M 2 are calculated on the basis of models [4] and are equal to M1 =150 and M 2 =90 for magnesium at T = 260 °C. Substituting these values into expression for Db* ( ε& ) , one can obtain that the ratio Db* ( ε& ) Db , grain-boundary diffusion enhancement factor, is equal to 35 at ε& =3×10-3 s-1. This is sufficient to accommodate sliding strain at grain triple junctions during grain boundary sliding, without which cavitation would immediately occur and cause sample fractured. Thus, taking into consideration of non-equilibrium state of grain boundaries, the achievement of superplasticity in microcrystalline ZK60 and AZ91 magnesium alloys at low temperatures of 200 - 300 °C can be reasonably understood.

6

Conclusions

1. ECAP can produce a homogeneous microcrystalline structure with non-equilibrium grain boundaries in high-strength magnesium alloys ZK60 and AZ91 alloys and to achieve low-temperature superplasticity at strain rate of 3 × 10-3 s-1: elongation to failure is equal to 810 % (260 °C) and 570 % (300 °C) in ZK60 and AZ91 alloys, respectively.

414 2. It is shown that the activation energies are 75 and 78 kJ/mol for the ECAP-processed ZK60 and AZ91 alloys, respectively. These values are much lower than 92 kJ/mol reported in literature for grain boundary diffusion in Mg. The enhanced grain boundary diffusion observed in the present study is caused by the fact that ECAP processed alloys possess non-equilibrium grain boundaries. This grain boundary diffusion enhancement greatly assists strain accommodation at grain triple junctions during superplastic deformation and promote low-temperature superplasticity.

7

References

[1] ASM Specialty Handbook Magnesium and Magnesium Alloys (Eds.: M. Avedisian and H. Baker), 1999, p. 314. [2] V. M. Segal, V. I. Kopylov, V. I. Reznikov, D. A. Pavlik, V. F. Malyshev, Processes of Plastic Structure Formation of Metals, Nauka i Tehnika, Minsk, 1994, p. 232. [3] V. I. Kopylov in Investigations and Applications of Severe Plastic Deformation (Eds.: T. C. Lowe and R. Z. Valiev), NATO ASI Series 3. High Technology, Kluwer Academic Publisher, Netherlands, 2000, 23 - 27. [4] V. N. Chuvil’deev, V. I. Kopylov, W. Zeiger, Annales de Chim. - Sci. des Matériaux 2002, 27, 55 - 64. [5] V. N. Chuvil’deev, V. I. Kopylov, M. Yu. Gryaznov, A. N. Sysoev, Doklady Physics 2003, 7, 21 - 24. [6] H. Watanabe, T. Mukai, K. Higashi, Scripta Materialia 1999, 40, 477 - 484. [7] H. Watanabe, T. Mukai, K. Ishikawa, K. Higashi, Scripta Materialia 2002, 46, 851 856. [8] M. Mabuchi, H. Iwasaki, K. Yanase, K. Higashi, Scripta Materialia 1997, 36, 681 686. [9] M. Mabuchi, K. Ameyama, H. Iwasaki, K. Higashi, Acta Materialia 1999, 47, 2047 2057. [10] K. Kubota, M. Mabuchi, K. Higashi, J. of Materials Science 1999, 34, 2255 - 2262. [11] S. Kamado, T. Ashie, H. Yamada, Materials Science Forum 2000, 350 - 351, 65 - 72. [12] T. Mukai, M. Yamanoi, H. Watanabe, K. Higashi, Scripta Materialia 2001, 45, 89 94. [13] M. Furukawa, Y. Iwahashi, Z. Horita, M. Nemoto, T. G. Langdon, Materials Sci. Eng. 1998, A257, 328 - 332. [14] T. G. Nieh, J. Wadsworth, O. D. Sherby, Superplasticity in metals and ceramics. Cambridge Univ. Press, 1997, p. 321. [15] H. J. Frost, M. F. Ashby, Deformation-Mechanisms Maps, Pergamon Press, Oxford, 1982, p. 166. [16] V. N. Chuvil’deev, A. V. Petryaev, Physics of Metals and Metallography 2000, 89, 122 - 126.

415

Structural Durability and Fatigue Life Assessment for Magnesium Diecast Specimen Karl Ludwig Kotte, Sven Werdin Dresden University of Technology, Institute of Solid State Mechanics

1

Introduction

Lightweight design asks for the use of magnesium in applications that are also relevant to functionality and safety. But the implementation of this material is restricted by the lack of knowledge about the operational strength of structures under various ambient conditions and the absence of design guidelines. The fatigue behaviour of two types of notched specimen made of the magnesium diecast alloys AZ91, AE42 and AM50 was investigated at constant amplitude and variable amplitude tests. The ambient conditions of these tests were laboratory air at room temperature, at 125 °C and a permanent influence of sodium-chloride solution at room temperature. Based on a substantial and consistent data set several variants of nominal-stress concept as well as local-strain approach were applied to draw conclusions to improve the reliability of lifetime estimation of components made of magnesium.

2

Specimen and test program

R 10

ø 10

ø 16

200 15

200

ø10

a)

185

26

6 mm thick

20

32

R

5

ø 16

Two types of notched specimen – one used in axial load and the other in bending tests – form the basis of this investigation, Figure 1. All parts were casted of the alloys AZ91, AE42 and AM50 using the same mould at the research foundry of the Audi AG. The mould contained an unnotched specimen too, Figure 1 c), to determine the cyclic behaviour of the alloys in incremental step tests and in cyclic stain controlled tests.

c)

R 40

b)

Figure 1. Specimen a) with notch factor Kt = 1.4 were used in axial load tests, specimen b) with notch factor Kt = 1.3 in bending tests and unnotched specimen c) to determine cyclic behaviour of the alloys

416 The specimens were tested in the as delivered condition. Only the ridge was deburred. Approximately 700 specimen were examined in 42 constant amplitude test series varying • between axial load and bending, • stress ratio R = -1 and R = 0, • alloys AZ91, AE42 and AM50, • ambient conditions with laboratory air at room temperature, 125 °C and a permanent influence of a 5% sodium-chloride solution at room temperature with and without a MAGOXID-COAT® to determine finite life fatigue strength. The tests were conducted in the range from 104 cycles to 2 ⋅ 107 cycles. Every single test series contained 8 tests at minimum. 403 variable amplitude tests were carried out in 44 test series under the same conditions mentioned above. Four synthetic load histories that differed in the kind of change of mean stress were used 0: stationary Gaussian random process without and with a constant mean stress ( R = 0), a load history that was developed from alternating mean stress every 1000 cycles and the Gaussian random process ( R = -1) and at last the overlay of the random process with a swelling mean stress every 100 cycles ( R = -0.5).

3

Results

3.1

Constant amplitude tests – slopes of S-N curves and mean stress sensitivity

All test series were analysed itself in the first step. The slopes k of S-N curves were fitted in the range from k = 3.1 to k = 6.4, Figure 2. AZ91 laboratory air corrosion MAGOXID + corrosion 125°C

axial load; R = -1 axial load; R = 0 bending; R = -1 bending; R = 0 generalized slope; R = -1 generalized slope; R = 0

AM50 laboratory air corrosion MAGOXID + corrosion 125°C AE42 laboratory air corrosion MAGOXID + corrosion 125°C 3.0

3.5

4.0

4.5 5.0 5.5 slope of S-N curve

6.0

6.5

Figure 2. Slopes of S-N curves for test series and generalized slopes

Already in the region of about 104 cycles to failure the ambient of 125 °C leads to a decrease of strength to 85% respectively to 80% under corrosive conditions. Test series at a stress ratio of R = 0 tended to result in smaller slopes then at R = -1. Only at 125 °C the slopes were all about the same. Generalizing these experiences all test series were ana-

417 lysed again with a target of slopes shown in Table 1. This procedure caused only a small increase of sizes of scatter TN defined by TN = N90%/N10%,

(1)

where N90% and N10% are the numbers of cycles at the probability of survival of 90% respectively 10%. But this generalization implicates better results in lifetime estimation. Table 1. Generalized slopes of S-N curves, averaged mean stress sensitivities

test conditions ambient condition stress ratio laboratory air 125 °C corrosion

-1 0 -1 0 -1 0

generalized slope of S-N curve AZ91 AE42 AM50 4.5 4.0 4.5 4.5 4.5 4.0

4.5 4.0 4.5 4.5 4.5 4.0

averaged mean stress sensitivity AZ91/AE42/AM50

6.0 5.0 6.0 6.0 6.0 5.0

0.55 0.30 0.60

10 000 20 000 50 000 100 000

basic test series additional test series

M=

100

M

ma x

=0 ,73

=

0

M/3

R

R = -1

50

0

0,4

8

stress amplitude Sa [MPa]

150

,5 R = -1

The mean stress sensitivity M is used to assess cycles with mean stress Sm deviating from that of S-N curve. The fatigue life of the researched alloys is very sensitive in terms of mean stresses. In addition caused by different slopes for swelling and alternating load the mean stress sensitivity depends on stress amplitude Sa respectively the number of cycles to failure, like shown in the Haigh-diagram in Figure 3. Averaged mean stress sensitivities can be taken from Table 1.

-100

-50

R=

0,5

0 mean stress Sm [MPa]

Figure 3. Haigh-diagram (AZ91, axial load, laboratory air)

50

100

418 3.2

Lifetime estimation following the nominal-stress concept

The aim of this investigation was the determination of the most reliable lifetime estimation procedure. The assessment is carried out by utilizing the ratio D = Nexperiment /Ncalculation (2) of numbers of cycles to failure. Service life was calculated for every single test in the data set varying elements of the nominal-stress concept. Mean value of damage sum D and size of scatter results from statistical analysis in the probability grid. The mean value D can be introduced as a correction factor into fatigue life estimation while the size of scatter TD is a measure of the unreliability of the calculation method. The size of scatter (3) TD = D10 % /D90 % is the ratio of cumulative damage sum at the confidence probability of 10 % and 90 %. It contains fractions of scattering of experiments and dispersion of calculation method. Several modifications of the Miner’s rule – Miner basic 0, Miner elementary, Miner modified 0, Miner consistent 0 and the modification by Liu/Zenner 0 – were tested. The load history was counted using rainflow algorithm. The mean stress transformation Sa trans = Sa + M ⋅ Sm (4) uses the Haigh-diagram 0. The kind of determination of mean stress sensitivity M was varied too. Either both original S-N curves (R = -1 and R = 0) from test series – variant a) – or the S-N curve with R = -1 and an averaged mean stress sensitivity – var. b) were applied. Utilizing S-N curves established with generalized slopes results in two more variants c) and d). Usually commercial software provides such combinations. The last method of transformation combines var. c) with an idea of Kotte 0, 0 that includes not only the mean stress itself into transformation procedure but also the reference to the mean stress of the whole load history – var. e). This method increases the influence to damage sum of load cycles with mean stresses below mean stress of the load history. The results of the most important calculations are shown in Figure 4. 0,38

Miner, elementary, a)

T = 13.9

0,33

Miner, elementary, b)

T = 16.8

Miner, elementary, c)

0,40

T = 10.6

Miner, elementary, d)

0,37

T = 12.6

0,14

Miner, basic, c) Miner, modified, c)

0,25

Miner, consistent, c)

0,25

T = 20.8 T = 12.5 T = 15.2 0,59

Zenner/Liu, c)

T = 8.8

0,80

Zenner/Liu, e) 0.01

0.1

1 damage sum D

10

T = 7.1 size of scatter 100

Figure 4. Statistical analysis of several variants of lifetime estimation using nominal-stress concept. Mean value, confidence band limited at a probability of 10 % and 90 % and ultimate extremal values were charted. Sample size: 403 tests. For explanation of calculation variants see text above.

419 3.2

Lifetime estimation following the local-strain approach

Basics of estimation of lifetime using the local-strain approach 0 are the determination of cyclic stress strain curves using the Ramberg-Osgood equation 0 (5) ε = σ/E+(σ/K')1/n' and strain-life curves fitted in the description of Manson/Coffin and Morrow 0 εa = εa,el + εa,pl = σf'/E ⋅ (2NA)b + εf' ⋅ (2NA)c . (6) Parameters were assorted in Table 2. Here the Neuber’s rule 0 and the damage parameter proposed by Smith/Watson/Topper 0 were applied. These standard assumptions are implemented in nearly all of the commercial software products using the local-strain approach. The local-strain approach considers lifetime until occurrence of a crack. Therefore crack initiation was identified using a potential drop instrument in many tests. The beginning of crack growth was averaged at 85% of cycles to failure. Table 2. Parameters to describe cyclic stress strain and strain S-N curve

alloy AZ91 AE42 AM50

temperature 20°C 125°C 20°C 125°C 20°C 125°C

E [MPa]

n'

39 180 35 470 36 580 33 370 37 240 33 790

0.3449 0.3907 0.4085 0.3889 0.3825 0.3557

K' [MPa]

σf' [MPa]

b

εf'

c

1 326 1 314 1 396 1 037 1 289 834

646.3 793.7 416.8 405.2 595.5 463.8

-0.1605 -0.2092 -0.1534 -0.1611 -0.1762 -0.1706

0.1240 0.2784 0.0523 0.0886 0.1316 0.1883

-0.4663 -0.5366 -0.3763 -0.4130 -0.4595 -0.4771

Statistical analysis of lifetime prediction with the basic procedure of local-strain approach results in small scatter, Figure 5. This is in contrast with nominal-stress concept. The tests under corrosive conditions were omitted in this analysis due to the absence of cyclic material data. local-strain approach 0.01

0,83 0.1

1 damage sum D

10

T = 8.9 size of scatter 100

Figure 5. Statistical analysis of calculation following the local-strain approach. Sample size: 254 tests (only tests in laboratory air and at 125 °C)

Some elements of estimation were modified to improve the reliability of prediction. But using the expanded damage parameter according to Bergmann 0 instead of the formulation of Smith/Watson/Topper or applying the advanced Neuber rule assuming elasticplastic material behaviour in the notch root had no downsizing effect to the size of scatter. According to its formulation the Smith/Watson/Topper damage parameter cannot cover the high mean stress sensitivity of magnesium. Some further developments are necessary to take this phenomena into account.

420

4

Conclusion

Calculation methods already successfully utilized to predict service life of steel or aluminium parts were applied successfully to estimate fatigue life of diecast magnesium parts with results that are comparable with the other material groups 0. Already using standard assumptions the local-strain approach achieves a high reliability of prediction. Applying the nominal-stress concept additional assumptions and procedures has to be implemented to achieve comparable results.

5

Acknowledgement

This paper presents results of a research project initiated and aided by the German Federation of Industrial Cooperative Research Associations "Otto von Guericke" (Arbeitsgemeinschaft industrieller Forschungsvereinigungen "Otto von Guericke" e.V.). The detailed report can be read in 0. Thanks to the AUDI AG for diecasting a large quantity of samples and to the AHCOberfächentechnik Holding GmbH for treating specimen with MAGOXID-COAT®.

6

References

[1]

C. Berger, K.-G. Eulitz, K. L. Kotte, H. Naundorf, W. Schütz, C. M. Sonsino, A. Wimmer, H. Zenner, International Journal of Fatigue 2002, 24, 603-625 J. W. Bergmann, Doctoral thesis, Darmstadt University of Technology, 1983 O. Buxbaum, K. L. Kotte, Erhöhung der Genauigkeit der Lebensdauervorhersage durch Berücksichtigung der Reihenfolge – Final report, Deutsche Forschungsgemeinschaft, 1994 L. F. Coffin, Journal of Applied Mechanics 1950, 17, 233-248 E. Haibach, Technische Mitteilung des LBF Nr. TM 50/70, 1970 E. Haibach, H. P. Lehrke, LBF-report No. FB-125, 1975 E. Haibach, Betriebsfestigkeit: Verfahren und Daten zur Bauteilberechnung, VDIVerlag, Düsseldorf, 1989 K. L. Kotte, state doctorate, Dresden University of Technology, 1987 K. L. Kotte, C. Berger, S. Werdin, H. Döcke, J. Senf, FKM Forschungsbericht Heft 263, 2001 A. M. Miner, Journal of Applied Mechanics 1945, 12, 159-164 H. Neuber: Journal of Applied Mechanics 1961, 28 , 544-550 W. Ramberg, W. R. Osgood, NACA, Technical Note No. 902, 1943 T. Seeger, A. Beste, VDI-Fortschrittsbericht Reihe 18 Heft 2 1977 K. N. Smith, P. Watson, T. H. Topper, Journal of Materials 1970, 5 (4), 767-778 H. Zenner, J. Liu, Konstruktion 1992, 44, 9-17

[2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15]

421

Microstructure/Property Relationships of Three Mg-RE-Zn-Zr Alloys L. Gill, G.W. Lorimer and P. Lyon*, Manchester Materials Science Centre *Magnesium Elektron Ltd, Manchester (UK)

1

Introduction

Nd has a high solid solubility in magnesium and the Mg-Nd system exhibits age hardenability. [1,2] Additions of Zn to Mg-Nd alloys suppress the binary precipitation sequence and forms a pseudo-binary system involving precipitation of a ternary phase. [3] Wilson et al [4] found that additions of 0.5wt% Zn to a Mg-3 wt% Nd alloy increased the time to reach peak hardness and delayed over-ageing at 200 °C. Additions of 1.35wt%Zn did not have a similar beneficial effect on the precipitation hardening reaction. Rokhlin [5] reported that increasing Zn content reduced Nd solubility and was associated with a decrease in strengthening during ageing. The present paper gives details of the ageing response of three Mg-RE-Zn-Zr alloys and the quench sensitivity of one.

2

Experimental

The alloys were produced at Magnesium Elektron’s experimental foundry in Manchester. The charges were made up from pure Mg, Zn, Nd and Gd. Zr was added in the form of a hardener. The alloys were sand cast into plate-moulds 20 cm x 20 cm x 2.5 cm. The alloy compositions are given in Table 1. Table 1. Alloy Compositions (weight %) and solidus temperatures

Alloy designation 8441 8442 8443

Solidus Temperature °C 530-535 525-530

Zn 0.007 0.6 1.3

Zr 0.48 0.46 0.47

Nd 2.7 2.8 2.8

Gd 0.35 0.29 0.28

The solidus temperatures of the three alloys, as determined from optical microscopy of solution treated and quenched samples, are given in Table 1. The higher Zn-containing alloy had the lower solidus temperature, a result in agreement with Rokhlin5. The three alloys were solution treated at 520 °C under 2%SF6 98%CO2, followed by a hot water quench at approximately 80 °C. Sample were aged at 150, 200, 250 and 300 °C in oil or salt baths. Tensile testing was carried out at room temperature and 150 °C in duplicate on samples of 32 mm gauge length and 5.64 mm diameter, on an Instron 1185 machine, at 0.5 mm per minute. After yielding the crosshead speed was increased to 20 mm per minute.

422 Jominy end quench testing was carried out on the 8442 alloy. The test procedure and specimen dimensions were those specified by BS4437. The experimental set-up was very similar to that used by Hart et al. [6] to determine the quench sensitivity of aluminium alloys. Test specimens were solution treated at 520 °C for 4 hours under an argon atmosphere. To measure the quench rate thermocouples were placed along the centre line of each test specimen at regular intervals from the quench plane. The cooling rates given are the average quench rates between 520 and 150 °C. For transmission electron microscopy (TEM) 3 mm diameter disk samples were electropolished in a solution of, 100 ml absolute ethanol, 15 ml 2-butoxyethanol, 3 ml acetic acid and 17.5 g sodium thiocyanate at a voltage of 22 V at 0 oC. Specimens were examined with a CM 200 TEM operating at 200 keV.

3

Results and Discussion

3.1

As-cast and solution treated alloys

Figure 1 shows optical micrographs of the alloys in the as-cast and solution treated condition. The microstructure of the as-cast alloys consists of approximately equiaxed dendrites of α-magnesium solid solution surrounded by a divorced eutectic of α-magnesium solid solution and Nd-rich second phase. The presence of a divorced eutectic in neodymium containing alloys is in agreement with work done by Leontis [1]. The microstructures are similar to that described by Mukhina [7] for the ML10 alloy (Mg-2.2-2.8wt%Nd-0.10.7wt%Zn-0.4-1wt%Zr), ‘a solid solution of Nd, Zn, and Zr in Mg and a eutectic of Mg solid solution and (MgZn)12Nd’. Following solution treatment, the volume fraction of the Nd-rich second phase decreased, Figure 1.

50μm 8441 As cast

8441 Solution treated

50μm 8442 As cast Figure 1.

50μm

50μm 8442 Solution treated

423

50μm 8443 As cast

50μm 8443 Solution treated

Figure 1. Alloys in as cast and solution treated state

The grain sizes of the three alloys in the as-cast and solution treated condition are all approximately 45 μm in diameter Solution treatment at 520 °C resulted in the formation of large, 0.1-1 μm, precipitates in the centre of the α-magnesium grains of the two Zn-containing alloys, Figure 1. TEM examination of these alloys, Figure 2, reveals a complex distribution of rod-shaped and equiaxed precipitates. Tikhova et al. [8] described the formation of similar precipitates during the solution treatment of the ML10 alloy and identified them as zirconium hydride and zinc zirconide as well as intermetallics of (MgZn)12Nd. It is possible that these large particles are the same precipitates, but they have not been positively identified.

Figure 2. TEM micrograph of large precipitates in the grain centre of the 8443 alloy (17500x) solution treated 8hrs 520°C + hot water quench

3.2

Ageing Response

The ageing response of the alloys at 150, 200, and 300 °C are shown in Figures 3, 4 and 5. The as-quenched hardness of alloy 8441 is slightly less than that of 8442, which, in turn, is less that that of 8443. At all temperatures increasing the Zn content of the alloy increases the time to reach peak hardness and reduces overageing of the alloy. There is an increase in peak hardness with increasing Zn content when ageing at 300 °C (Figure 5). Peak hardness values at 150 and 200 °C are similar for the three alloys. In comparing this result to work done by Wilson et al. [4] on similar alloys, it can be seen that in their alloys, Mg-3wt% Nd, Mg-3wt% Nd-0.5wt% Zn, and Mg-3wt% Nd-1.35wt% Zn, no such

424 trend was observed. The high Zn alloy over-aged and reached peak hardness at the same rate as Zn-free alloy. This effect was deemed to be due to, ‘orientation of precipitate plates (contained in the Mg-3wt%Nd-1.35wt%Zn alloy) that provide the least barrier to dislocation motion.’ Ageing at 150°C

80 75

8441 150°C 8442 150°C 8443 150°C

VH

70 65 60 55 50 10

10

100 Log time hours

1000

10000

Figure 3. Hardness versus time at 150 °C Ageing at 200°C

80 75 Vh

70 65 60 55 50 10

10

100 Log time hours

1000

10000

Figure 4. Hardness versus time at 200 °C Ageing at 300°C

70

VH

65 60 55 50 45 0.01 0

0.1

1 10 Log time Hours

100

1000

Figure 5. Hardness versus time at 300 °C

3.3

Tensile Properties

Table 2 shows the UTS and elongation of the alloys at room temperature and 150 °C. The heat treatment given to the alloys were 8 hrs at 520 °C + hot water quench + 24 hours at 200 °C + cold water quench. From Table 2 it can be seen that there is a decrease in UTS

425 with increasing Zn content at room temperature and 150 °C. Rokhlin [5] gave UTS data for the Russian alloy ML10 as maximum UTS for T6 condition of 226 Mpa following an air quench and 235 MPa following a water quench. This is similar to the results shown in Table 2. Alloy 8441 has slightly better tensile properties than 8442 and 8443 at both room and elevated temperatures. Table 2. UTS and elongation of alloys at 20 and 150 °C

Alloy 8441 8442 8443 3.4

UTS 20 °C MPa 262 260 250

UTS 150 °C MPa 255 225 223

Elongation at failure % 20 °C

Elongation at failure % 150 °C

5 4 5

8 7 12

Jominy End Quench Results

Figure 6 shows the effect of quench rate on the ageing response of the 8442 alloy at 200 °C. It can be seen that after solution treatment the slowest quench rate (3 °C/min) has the highest hardness value and the fastest quench rate (43 °C/min) the lowest hardness. At peak hardness this relationship is inverted and the fastest quench rate has the highest hardness. At the slower quench rate large, heterogeneously-nucleated precipitates form during cooling and there is less solute in solid solution to form strengthening precipitates during subsequent ageing. With a faster quench the heterogeneously-nucleated precipitates do not have time to form and initially the hardness of the alloy is lower. Upon subsequent ageing the hardness increase is larger as there is more solute to form strengthening precipitates. 85

80

43°C/m in

Vickers Hardness

75

13°C/m in 70

8°C/m in 6°C/m in

65

3°C/m in 60

4°C/m in 3°C/m in

55

50 0.1

1

10

100

1000

Log Tim e (hours)

Figure 6. Ageing response at 200 °C for alloy 8442 quenched at rates between 43 °C/min and 3 °C/min.

4

Conclusions

The ageing response of the Mg-Nd-Gd-Zr alloy was changed by adding Zn. Increasing the Zn concentration increased the peak hardness of the alloy at 300 °C. At all temperatures increasing the Zn content increased the time to reach peak hardness and delayed the onset of overageing.

426 Additions of Zn to a Mg-2.7wt% Nd-0.3wt% Gd-0.5wt% Zr resulted in the precipitation of large (0.1-1 μm) precipitates during solution treatment at 520 °C. The UTS of the alloys in the T6 condition decreased slightly with increasing Zn content while the elongation was similar. Testing at 150 °C resulted in a decrease in the UTS of the all the alloys and increasing the Zn content caused a greater reduction in UTS. Elongation increased at 150 °C for all three alloys. The Jominy end quench test has been successfully used to determine the quench sensitivity of the Mg-2.7wt% Nd-0.3wt% Gd-0.6wt% Zn-0.5wt% Zr alloy.

5

References

[1] T.E. Leontis, Transactions AIME, 1949, 185, 968-983 [2] E.F. Emley, ‘Principles of magnesium technology,’ First edition, 1966, Pergamon press, London, (Now out of print, photocopy available from Elsevier science LTD, Oxford) [3] P. A. Nuttall, T. J. Pike, B. Noble, Metallography, 1980, 13, 3-20 [4] R. Wilson, C. J. Bettles, B. C. Muddle, J. F. Nie, Conference proceedings: Magnesium alloys 2003 PSTAM 2003, Grand Cube Osaka 26-30 January 2003, Trans Tech [5] L.L. Rokhlin, ‘Magnesium alloys containing rare earth metals,’ 1st edition, Taylor and Francis, London [6] W. G. J. Hart, H. J. Kolkman, L. Schra, National aerospace laboratory NLR, The Netherlands, NLR TR 80102 U [7] I. Y. Mukhina, V. M. Lebedev, K. H. Kim, D. K. Kim, Metal science and heat treatment, 39, 5-6, 1997, 202-206, Translated from Metallovedenie i termicheskaya obrabotka metallov 5, 1997, 21-25, Plenum Publishing [8] N. M. Tikhova, V. A. Blokhina, A. P. Antipova, T. P. Vasil’eva, G. G. Solov’eva, All union institute of aviation materials Translated from Metallovedenie i termicheskaya obrabotka metallov 10, 1971, 73-76, Plenum Publishing

6

Acknowledgements

The authors acknowledge financial support from EPSRC and Magnesium Elektron Ltd.

427

Influence of Extrusion Methods on Microstructure and Mechanical Properties of AZ31 Magnesium Alloy J. Dzwonczyk, N. Hort, K. U. Kainer GKSS Forschungszentrum GmbH, Geesthacht, Germany

1

Introduction

Extrusion – as a thermo-mechanical process has a long industrial history. In the past 50 years its economic importance has increased, predominantly as the result of significant technological advances that have drawn on extensive practical experience and on numerous fundamental investigations into the extrusion process, tooling and metal flow. For magnesium the extrusion process is applied with continuously rising success. There are no specific restrictions or limitations in magnesium extrusion different from any other materials. The process results in a very large range of possible geometries (bars, rods, shapes and tubes) in combination with maximum structural efficiency and minimum weight and costs [1,2]. Substantial progress has been made in the development of standard extrusion techniques. To the basic methods belong direct (forward extrusion) and indirect (backward extrusion). Another kind - hydrostatic extrusion - has presently started to play an important role in the processing of magnesium alloys [3,4]. In accordance to the scope of this study, it is useful to compare the main differences between the mentioned extrusions methods, which are considered in presented investigations. During direct extrusion a metal billet is pressed through a die by a ram while during indirect extrusion the die is placed into a hollow ram, which is pressed into the billet. In hydrostatic extrusion the pressure is applied to the billet by a pressure medium, which fulfils the space between the container and the billet. All three methods are presented schematically in Figure 1.

Figure 1. a) direct extrusion b) indirect extrusion c) hydrostatic extrusion.

Within the scope of this study investigations are aimed at a comparison between the influence of the different extrusion methods on materials properties at a constant extrusion temperature. Additionally the influence of different extrusion temperatures on the material had been analysed as extrusion temperature is a very important process parameter.

428

2

Experimental Procedures

AZ31 billets were subjected to direct and indirect extrusions at 275 °C, 300 °C and 350 °C. The extrusion ratio for direct and indirect methods was 23:1 and resulted in 20mm diameter round profiles. Hydrostatic extrusion trials were performed at 250 °C and 300 °C. In the accordance to the requirement of the hydrostatic method, the billets were machined from one side of the billet to the truncated cone shape. This operation allowed for tighter contact between the billet and the walls of the dies. Castor oil was used as the extrusion medium. The extrusion ratio for each hydrostatic trial was 26:1 and resulted in 15 mm in diameter of the extruded profiles. Specimens for microstructural analysis were chosen from the central part of extruded rods. The surfaces of the samples were prepared using a graded sequence of abrasives, starting with a coarse abrasive 500 mesh and then progressing up to 2400 grit. Applied polishing procedure was carried out employing 3, 1 and 0.25 μm diamond suspensions. The etching mixture was based on picric acid and ethanol. The specimens were etched in the solution for 20 sec. and then subjected to microstructural observations. The specimens for uniaxial tension tests were machined according to the standard DIN EN 50125 with a gauge length of 25 mm, diameter of 5 mm and fine threaded screw heads. Compression testing was done on the specimens with a length of 15 mm and diameter of 10mm and prepared in accordance to DIN EN 50106. For testing the specimens were chosen from the central part of extruded bars and then numbered according to the position by marking the direction of the extrusion. Both tests were performed with a de-3 -1 formation rate of 3.3 x 10 s , at room temperature. The specimens were also subjected to hardness measurements. In accordance to DIN EN 50125 the Vickers hardness tests was carried out on the material. For hardness tests the load of 9.8 N was submitted for 10 sec. For the evaluation of the mechanical properties standard equations have been applied [5].

3

Results and Discussion

3.1

Microstructural observations

Analysing the microstructures of AZ31 alloy extruded by direct, indirect and hydrostatic method at 300°C, it was seen that, although the direct and indirect methods vary from the process character itself, the grain structure shows no significant differences between each other – see Figure 2.

Figure 2. Microstructure of AZ31 processed at 300°C using a) direct, b) indirect and c) hydrostatic extrusion method.

429 The results of the average grain size analysis, in accordance to the effect of extrusion methods on the material, are presented in Table 1. It can be clearly distinguished that apart from the comparable effects of direct and indirect method on the material’s microstructure, hydrostatic extrusion changes the material’s character considerably. Although some percent of elongated grains were observed, the material extruded hydrostatically exhibits a very fine-grain structure compared to the other methods. This result indicates that by hydrostatic method of extrusion is possible to obtain material with a fine structure and in accordance to Hall-Petch relation – better mechanical properties can be expected [5,6]. It is necessary to mention that at the same time the amount of elongated grains in hydrostatically extruded materials was larger than that in extruded via direct and indirect method. Table 1. The influence of different extrusion methods and extrusion temperatures on average grain size of AZ31.

Extrusion temperarture

Direct extrusion

Indirect extrusion

Hydrostatic extrusion

250 °C 275 °C 300 °C 350 °C

11 ± 0,90 μm 12 ± 1,68 μm 16 ± 1,25 μm

11 ± 0,90 μm 13 ± 1,47 μm 16 ± 1,74 μm

2 ± 0,96 μm 3 ± 1,13 μm -

The observed behaviour seems to be an effect of strongly reduced frictional heat during the hydrostatic processing, where there is no additional heat reaction between the container and the billet and very small one between the material and the die. This could directly influence the dynamic recrystallization (DRX) state during processing. Some of the elongated grains include the traces of sub-grains formation, which indicate the DRX has occurred in the developing state. Additionally, the explanation can be hinted by the fact, that the material before the performance to hydrostatic extrusion was machined from the one side of the billet to get a better fit between the billet and the die shape. This operation automatically allows a reduction of the friction forces on the dies, which usually strongly affect the final recrystallization state of the extruded product. The results of the influence of the different temperatures applied to the extrusion trials on average grain size of extruded material are clearly indicated in Table 1. It can be concluded that the increase of the extrusion temperature in the range of 275 °C to 350 °C has a slight, but remarkable influence in the case of the direct and indirect methods: with an increase of the extrusion temperature, the average grain size also increases – from around 11 μm at 275 °C to 16 μm at 350 °C, while for the hydrostatic method with increasing temperature (from 250 °C to 300 °C) the average grain size is practically constant (2 μm and 3 μm for 250 °C and 300 °C, respectively). 3.2

Mechanical Testing

Hardness measurements applied to the investigated alloy pointed out comparable results for direct and indirect extrusions, while in the case of hydrostatic extrusion the Vickers hardness increases considerably. This result is in good agreement with microstructural investigations. The amount of the elongated grains, which were detected in the microstructure, seems to have a negligible effect on the materials hardness. Analysing the effect of different extrusion temperatures on this property there was found that the extrusion temperature in the applied range does not significantly influence the material properties

430 and there is only a slight decreasing tendency of hardness in direct and indirect extruded material was visible. The comparison between the hardness investigations on AZ31 is presented in a Figure 3. 100

DIRECT

HYDROSTATIC

compression

300

Strain [%]

Vickers Hardness

400

INDIRECT

80

60

tension 200

40 100

20 0 0

5

10

250°C

275°C

300°C

15

Stress [MPa]

0

20

25

350°C

Extrusion Temperature

Figure 3. The influence of different extrusion methods and extrusion temperatures on hardness of AZ31. 600

Tensile Strength Strain to Fracture (Tension) 0.2% Yield Strenght (Compression)

500

Figure 4. The asymmetry between material flow during tensile and compression test (AZ31, direct extruded at 300 °C).

0.2% Yield Strength (Tension) Compressive Strength Strain to Fracture (Compression) 420

Stress [MPa]

400

368

363 319

300

200

269

268

268 183

240

179 114

113

100 23

11

23

11

24

0 DIRECT

INDIRECT

8

HYDROSTATIC

Extrusion Method

Figure 5. The influence of different extrusion methods on mechanical properties of AZ31 extruded at 300 °C.

The observed significant difference between mechanical properties after uniaxial tension and compression results from strong anisotropy, which is a pattern feature of hcp metals. After mechanical testing performed on extruded AZ31 the material flow curves exhibit considerable asymmetry between each other, which can be clearly seen in a Figure 4. This result directly indicates the connection with the texture of the investigated material [7, 8]. Especially in the case of the extrusion of Mg-based materials the basal planes tend to be strongly aligned parallel to the axis of the rod. This fact strongly reflects on the behaviour of the material under tensile and compression stresses. The tensile stresses applied perpendicularly to basal plane support twin formation, while for stresses parallel to basal planes twins are possible only in compression. Taking under discussion the influence of different extrusion methods on the material properties it is important to point out that via direct and indirect extrusion at 300 °C the similarity in the results can be accomplished but with a great difference between yield stress and ultimate stress both in compression and tensile mode. The material extruded via hydrostatic method achieved

431 much better mechanical properties with close values of tensile yield and ultimate tensile stress but with a great difference in yield and ultimate compression under compressive load, again with a great agreement to the microstructural observations and the results of the hardness measurement. Another agreement correlates with Hall-Petch relation between the grain size and yield strength of the material. Tensile and compression results for the materials extruded at 300 °C are presented in a Figure 5. All three materials indicate a properties’ dependency on the grains size. In hydrostatic extruded material, where the grain size is in the range of 2 μm, the compressive 0.2% yield strength reached the value higher than 100% of the yield strength of the material extruded directly and indirectly. While the yield stress and the grain size are in good aggreement with the Hall-Petch predictions this was not the case in compressive testing. Here the reduction of the grain size leads to a decrease of compressive yield stress (Fig. 6). For comparison between the mechanical properties of the material subjected to the extrusion with changed extrusion temperatures the direct extruded material was chosen. The measured properties are presented in Figure 7. 300

250 YS, tension

Stress [MPa]

200

150

100 YS, compression 50

0 0,25

0,30

0,35

0,40

d

-1/2

0,45

0,50

0,55

0,60

Figure 6. The relationship of .compressive/tensile yield stress and grain size 500 450

Tensile Strength Strain to Fracture (Tension) 0.2% YS (Compression)

400

375

0.2% YS (Tension) Compressive Strength Strain to Fracture (Compression) 368

365

Stress [MPa]

350 300

269

269

266

250 200

192

183

180

150

117

114

101

100 50

23

11

23

11

20

10

0 275°C

300°C

350°C

Extrusion Temperature

Figure 7. The mechanical properties of AZ31 (direct extruded) with different extrusion temperatures.

432 The results indicate that the changes in the material among varied extrusion temperatures in the range from 275 °C to 350 °C are very small. However, the general tendency was that increasing the extrusion temperature in the investigated range results in negligible change in ultimate tensile strength of the material. At the same time the tensile 0.2% yield strength was increased approximately 10 MPa and the strain to fracture decreased about 3%. The compression results indicated a similar tendency for compressive strength, but in the case of 0.2% yield strength the inclination was opposite to the tensile tests and the decrease of the YS value could be observe (from 117 MPa to 101 MPa). The same trend was observed in the material extruded via indirect method among different extrusion temperatures. In the case of the material extruded via hydrostatic method there was no differences in the mechanical properties among the range of investigated extrusion parameter.

4

Conclusions

The investigation program subjected to the AZ31 magnesium wrought alloy extruded via direct, indirect and hydrostatic method grand the following conclusions: In accordance to different extrusion methods (AZ31 extruded at 300 °C): The microstructural observations indicate large structure similarities for direct and indirect methods, while the hydrostatic extrusion considerably influences the structure by significant grain refinement. The mechanical properties of the material extruded directly and indirectly are also similar in comparison to hydrostatically extruded trials. The hydrostatic method allows obtaining a very fine microstructure and very good mechanical properties of the extruded material. In accordance to different extrusion temperatures: Within the changes of the extrusion temperature in the range of 275 °C to 350 °C for direct and indirect extrusion and in the range of 250 °C to 300 °C for hydrostatic one the main indications are that for direct and indirect method the small changes in properties were observed, while for hydrostatic extrusion the change of 50 °C seems to have a unique influence on the microstructure and mechanical properties of the material. In the same way the direct and indirect methods give the impression that they are more sensitive to the changes of the extrusion temperature.

5

References

[1] F. E. Katrak, et al. in Magnesium Technology 2000 (Ed. H. I. Kaplan, et al.), 2000, p. 351-333 [2] D. Brungs, H. Gers, A. Mertz, Honsel GmbH & Co. KG, 2001 [3] H. J. McQueen, et al. in Magnesium Technology 2000 (Ed. H. I. Kaplan, et al.), 2000, p. 359-360 [4] K. B. Mueller in Magnesium Technology 2002 (Ed. H. I. Kaplan, et al.), 2002, p. 187-192 [5] G. D. Dieter, Mechanical metallurgy (third ed.), 1986, p. 275-290 [6] F. J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon Press, 1995, p. 52-53 [7] M. Hilpert, L. Wagner in Magnesium Technology 2000, (Ed. H. I. Kaplan, et al.), 2000, p. 375-381 [8] I. L. Dillamore, et al., Texture 1972, Ed. Engl., vol.1, p. 17-29

433

Anisotropic Mechanical Properties of Extruded Magnesium Alloy AZ31 C.H.J. Davies, F. Xiong, and M. Battaini Cooperative Research Centre for Cast Metals Manufacturing (CAST), School of Physics and Materials Engineering, Monash University, Victoria, Australia, 3800

Introduction If predictions [1] are to be believed, the use of wrought magnesium is set to increase dramatically over the next five to ten years as auto manufacturers and others seek to lightweight components as diverse as cars and luggage. However, this increased use will rely in part on improved understanding of the forming behaviour of wrought magnesium alloys, both at elevated and room temperatures. This paper examines the combined effects of test orientation and c-axis orientation on yield anisotropy at room temperature.

Experimental Procedure Extruded AZ31 magnesium alloy was supplied as rectangular cross section bar with a thickness of 6mm and as round rod with diameter 28mm. Tensile specimens were machined from the rectangular bar at an orientation parallel to the extrusion direction and perpendicular to the extrusion direction; for the rod, tensile samples were taken only in the extrusion direction. Plane strain compression (PSC) was conducted for several sample orientations (Table 1); for the rod, the PSC samples were taken as close as possible to the centerline of the rod. All specimens were deformed at room temperature in an Instron mechanical testing machine at a true strain rate of 0.01 s-1. Table 1. Relationship between plane strain compression sample axes and principal directions in the extrudate for bar and rod samples.

Compression Bar, orientation 1 Bar, orientation 2 Rod, orientation 1 Rod, orientation 2

Extrusion Transverse Radial Radial

Sample axes Constraint Normal Normal Extrusion Radial

Extension Transverse Extrusion Radial Extrusion

Texture was measured using an MMA texture goniometer, using Cu Kα radiation operated at 40 kV and 25 mA; the machine was fitted with polycapillary optics. A reference random sample was cold isostatically pressed from powder of the alloy.

434

Model In an attempt to explain variations in proof stresses for the various testing orientations and two starting textures, a simple upper bound model was developed based on an analysis of the Schmid factor variation for the different possible slip / twinning systems operating. The basis of the model is as follows [2]. The shear stress acting on a slip plane is given in tensor notation by the well known expression τ = bi n jσ ij

(1)

where the slip direction is b, and the slip plane normal, n. This equation may be transformed to τ = bi ( Tik T jlσσ kl′′ ) n j

(2)

where cos φ2 cos φ1 − cos Φ sin φ2 sin φ1 Tij = − sin φ2 cos φ1 − cos Φ cos φ2 sin φ1 sin Φ sin φ1

cos φ2 sin φ1 + cos Φ sin φ2 cos φ1 sin Φ sin φ2 − sin φ2 sin φ1 − cos Φ cos φ2 cos φ1 sin Φ cos φ2 (3) − sin Φ cos φ1 cos Φ

where the crystal orientation (Euler angles) is expressed in Bunge notation, and σ’’ is the stress matrix in the sample reference frame; σ is the applied stress. For tension in the extrusion direction, the only non-zero term is σ11, whereas for tensile stress in the transverse direction σ22 is non-zero. In this case, the Schmid factor is given by the product m = bi ( Tik T jlσ kl′′ ) n j

(4)

Schmid factors for the various slip (twinning) systems (Table 2) may then be determined for the range of Euler angles defined by the initial texture of the samples. Although there is dispute as to whether twinning can be characterized by a critical resolved shear stress, we make the assumption that a Schmid factor can be calculated for the known twinning plane and shear direction. Five possible deformation systems were included in the model (Table 2) and the relative critical resolved shear stresses (CRSS) of each were gathered from literature [3,4]. For both the bar and the rod, uniaxial tension in the extrusion direction was examined. For the bar only uniaxial tension in the transverse direction was also evaluated. Deformation of the bar was examined in plane strain compression in the extrusion direction, and in the transverse direction. For the rod, deformation was examined in plane strain compression such that the constraint was either in the extrusion direction or in the radial direction. Using the method and assumptions above, the ratio of the yield stress between two test orientations may be calculated for any crystal orientation for any applied stress state. Because of the uncertainty in the absolute value of CRSS for each system, yield stress ratios were calculated.

435 Table 2. Deformation systems modeled. (Tension twinning was allowed to operate in both tension and compression, whereas compression twinning was allowed to operate only in compression.)

System

Slip / twinning plane

Slip / twinning direction

Relative CRSS

Basal Prismatic, a Pyramidal, c + a Tension twinning Compression twinning

{0001} {1 1 00} {11 2 2} {10 1 2} {10 1 1}

<11 2 0> <11 2 0> <11 2 3> <10 1 1 > <10 1 2 >

1 30 30 4 12

Results The initial textures of the two extruded products were as expected. Both had a strong intensity of c axes normal to the direction of extrusion, but because of radial symmetry, the extruded rod had more of a spread in the ‘transverse’ (ie, radial) direction (Figure 1).

(a)

(b)

Figure 1. {0001} pole figures for (a) extruded bar, and (b) extruded rod.

In common with the findings of others, when tested in tension the proof stress for the bar parallel to the extrusion direction (ED) was greater than the proof stress perpendicular to the extrusion direction for all strain rates. The form of the stress strain curve was also different, with greater work hardening in the perpendicular orientation (Figure 2). Plane strain compression of the bar showed little difference in proof stress between the two test orientations, but the expected tension/compression yield anisotropy was evident (compare Figures 2 and 3). In contrast, the rod showed considerable difference between the two orientations in plane strain compression, with the test constrained in the radial direction having a higher proof stress than the test constrained in the extrusion direction (Figure 4). The tensile proof stress in the extrusion direction was, again as expected, higher than both, but not significantly different from the ED proof stress of the bar (Figure 2).

436

Equivalent true stress (MPa)

350 300

ED bar ED rod

250 200

TD bar

150 100 50 0 0

0.05

0.1

0.15

0.2

0.25

Equivalent true strain

Figure 2: Tensile stress strain curves for extrusions tested in orientations as labelled.

Equivalent true stress (MPa)

600 500 400

OR1 OR2

300 200 100 0 0

0.05

0.1

0.15

0.2

Equivalent true strain

Figure 3. Plane strain compression of bar extrudate. Sample OR1 is compressed in the extrusion direction; sample OR2 is compressed in the transverse direction.

We calculated the ratio of yield stress in the extrusion direction to yield stress in the transverse direction, in both tension and compression, for two c axis orientations. In essence the orientations chosen represent points in the continuum of orientations described by the experimental pole figure (Figure 1a). For a third c-axis orientation, we calculated the ratio of yield stress for ED constraint to the yield stress for TD constraint. This third orientation was chosen as an archetype for the extremes of the pole figure for the rod extrudate. For the first two orientations, we also compare the yield stress in tension to the yield stress in plane strain compression. Predictions of the ratio of yield stresses for three c axis orientations (Table 3) correlate well with our experimental observations, which for the conditions in Table 3 were σ90/σ0 = 0.77 in tension, and 1.0 in PSC for the bar and for the rod 0.6, whereas σtension/σPSC was 2.05 for the bar.

437 Table 3. Yield stress ratio predictions for three c axis orientations.

Angle of c axis to ND σ90/σ0 Tension 15o 0.76 45o 0.72 85o

σ90/σ0 PSC 0.94 1.12 0.62

σtension/σPSC in ED 1.23 1.49 -

Equivalent true stress (MPa)

350 300 250

OR1

200 OR2

150 100 50 0 0

0.02

0.04

0.06

0.08

0.1

0.12

Equivalent true strain

Figure 4. Plane strain compression of bar extrudate. Sample OR1 is constrained in the radial direction; sample OR2 is constrained in the extrusion direction.

Discussion One cannot easily use the Schmid factor model to evaluate the complete range of texture components (ie using the ODF as a starting point), but it is principally the orientation of the c axis that is of interest. In this case we have selected three c axis orientations (relative to the ND) to model the extrudate. Given this, it is remarkable that the predicted σ90/σ0 ratios for tension and plane strain compression are so close to the experimental ratios, and we regard the correlation as somewhat fortuitous as the system is not closed (ie, there is more than one combination of CRSS values which can be coupled with the calculated Schmid factors to give the ‘correct’ ratios). However, we can make qualitative observations based on the predictions. Although it is principally basal slip that is the difference between our ED and TD predictions in tension, we also observe that tensile twinning is activated, and slip is more likely to become active in the TD. The difference in twinning activation is in accord with previous observations made on the same material [5]. In tension in the extrusion direction, very little twinning is observed, whereas tension in the transverse direction results in a large fraction of twins, and it is this that accounts for the different work hardening behavior of the two orientations [5]. The experimental observation that in the bar the two test orientations result in very similar PSC proof stresses (Figure 3) is explained in the model by the activation of basal slip in both orientations under the plane strain compression state, where both σ11 and σ22 take a non zero value.

438 We are able to predict the behaviour of the rod because the c axes are distributed in the radial direction and it is only our pole figure representation which forces the texture into a Cartesian reference frame (ED-TD-ND). Thus the single c-axis orientation chosen may realistically represent the rod texture. For this orientation the Schmid factor for basal slip is very low, whereas tensile twinning is active for both constraint in the ED and in the TD because in both of these test conditions c axis extension will occur.

Conclusions A simple model has been used to account for the yielding behaviour of extruded AZ31 bar and rod tested under tension and plane strain compression a range of orientations relative to the extrusion direction. Ratios of the yield stresses between tension and plane strain o o compression and between 0 and 90 relative to the extrusion direction are predicted quite well using just a few c-axis orientations and five possible deformation systems.

Acknowledgements The CRC for Cast Metals Manufacturing (CAST) is supported in part by the Australian government’s Cooperative Research Centres scheme.

References [1] H. Friedrich, S. Schumann, Proc. IMA Conference 2001, 8-14. [2] C.N. Reid, Deformation Geometry for Materials Scientists, Pergamon, 1973. [3] M.R. Barnett, in Proc “Thermomechanical processing: microstructure, mechanics and control” held Sheffield, UK, June 23-26 2002; in press. [4] S.R. Agnew, M.H. Yoo, C.N. Tome, Acta mater., 2001, 49, 4227-4289. [5] F. Xiong, and C.H.J. Davies, in Thermec 2003 (Ed.: T. Chandra, J.M. Torralba, T. Sakai), Materials Science Forum, 2003, vols 426-432 Trans Tech Publications, pp3605-3610.

439

Mechanical Properties of Hot-rolled Magnesium AZ31 Sheets J. Wendt, V. Šupik, P. Zhang and J. Zhang Chair of Physical Metallurgy and Materials Technology, Technical University of Brandenburg at Cottbus, Cottbus, Germany

Abstract Sheets of magnesium alloys are one important option for the mass reduction in car manufacturing. Within the project “MIA” – Magnesium in Car Manufacturing sheets of the magnesium wrought alloy AZ31 were produced using a hot rolling mill. The mechanical properties of these sheets were investigated for various thicknesses of the sheet material. All these results including the surface properties like surface roughness and the microstructure of the specimens were compared with commercially available magnesium AZ31 sheets.

1

Introduction

Fuel consumption is one of the most important aims for the development of new cars. A very effective way is the reduction of mass. The last years showed that structures of aluminum alloys can replace steel components. But the development of new safety technologies in the cars or components like air condition annuls this advantage. New, lighter materials are needed. The application of magnesium alloys, whose density is only roughly 25% that of steel and 66% that of aluminum is expected to substantially increase in this decade. The most applications of these alloys are concentrated on using cast parts. However, the use of sheet materials will help further to reduce the mass of new cars. The big problem for producing of magnesium sheets is the well known poor deformation behaviour at room temperature [1–4].

2

Experimental

The wrought magnesium alloy AZ31 (nominal composition in weight percent: 3 Al, 0.8 Zn, 0.2 Mn, balance: Mg) was received as extruded rectangular bar stocks from Otto Fuchs Metallwerke Meinerzhagen, Germany and Honsel Profilprodukte Soest, Germany. The as-extruded material was after preheating rolled on a pilot hot-rolling mill (Figure 1) at temperatures of the rolls of 300 °C with different rolling parameters such as preheating temperature and time and pass deformation degree [5–7]. The dimensions of the rolling materials were 450 mm × 200 mm × 20 mm. Rolling was done in 11 passes. The thickness of the sheets was 1 mm. The magnesium sheets had a width of 600 to 650 mm and a length of 1000 to 1200 mm. The rolling direction was perpendicular to the extrusion direction. Crystallographic textures of the sheets were measured by X-ray diffraction using Nifiltered Cu-Kα radiation. Surface roughness measurements in rolling direction (RD) and transverse direction (TD) were done. Tensile tests were performed on flat specimens for

440 the hot-rolled sheets and on round specimen having gage lengths of 20 mm for the as received material. The initial strain rate was 8.3 × 10–4 s–1. The specimens were taken from different directions in the sheet. For each direction five specimens were tested. Fracture surfaces were observed with SEM. For comparison tensile tests were performed with other sheets of the magnesium alloy AZ31 from Spectrulite, USA and a producer from Russia. For these two sheets the manufacturing history concerning rolling and finishing is unknown.

Figure 1. Isotherm hot rolling mill

3

Results and Discussion

The microstructure of the cross section of the as received materials is shown in Figure 2a. The average grain size in the middle of the cross section is less than 100 μm. After rolling a very fine grain structure was observed. The average grain size directly after cooling in air is 8 μm (Figure 2b). The microstructure of the two other commercially available AZ31 sheets are shown in Figure 3. The average grain size for both sheets is less than 10 μm.

a) extruded material Figure 2. Microstructure of AZ31 material

b) rolled material (total deformation degree: temperature of the rols: 300 °C

441

a) Spectrulite sheet

b) Russian sheet

Figure 3. Microstructure of AZ31 sheets

A typical pole figure of the extruded material is illustraded in Figure 4a. Most of the hexagonal magnesium grains in the as-fabricated extrusion of AZ31 are oriented with the basal planes parallel to the extrusion direction in the extrusion plane. For the as-rolled condition an orientation of the hexagonal unit cell with the basal plane parallel to the rolling plane can be observed. As it is shown in Figure 4b, one can expect an isotropic distribution of the mechanical properties inside the rolling plane.

a) as-extruded

b) as-rolled

Figure 4. (0002) pole figures of AZ31 for the various conditions

The measured surface roughness of one hot-rolled sheet is shown in Table 1. Investigations were performed after surface cleaning with 5% HNO3 in alcoholic solution. Values of Ra are nearly constant with 1.3 μm for the top and bottom of the sheet in both, rolling (RD) and transverse (TD) direction. Table 1. Surface roughness of the hot-rolled and cleaned AZ31 sheet

side top bottom top bottom

direction RD RD TD TD

Ra [μm] 1.3 1.2 1.3 1.4

Rz [μm] 6.5 6.9 6.7 7.9

Ry [μm] 7.6 9.4 8.5 9.7

442 Tensile tests at the extruded material were done in the longitudinal and in the long transverse direction of the bar stock. The directions of the speciemens taken from the sheets for tensile tests were rolling direction (0°), transverse direction (90°) and 45° as it is shown in Figure 5.

Figure 5. Testing directions in extruded and rolled material

The result of the tensile tests are collected in Table 2. The values for the as extruded material are strongly depending on the direction. For example the yield strength of the long transverse directions is only 66 % of that of the longitudinal direction.The difference in the ultimate tensile strength is less than 35 MPa. This behaviour can be explained by the crystallographic texture which is shown above. For the rolled material there are only small differences in mechanical properties for the different directions. This is corresponding to the (0002) pole figure shown in Figure 4. For example the maximum difference in yield strength is only 5 MPa between 0°-direction and 45°-direction. In addition, for the values of elongation there is not significant difference. Comparing to the as-extruded material the vield strength is only 10 to 15 MPa lower than that in the longitudinal direction of the bar stock. The ultimate tensile stength is in the same value as for the long transverse direction in the extrusion and the ultimate elongation is nearly twice as high as in the extruded material. In comparison with the Spectrulite and the Russian sheets, the yield stength is higher than that of the other sheets. For the ultimate tensile strength the Spectrulite sheet shows the highest values. The elongation values of the other sheets are higher than values of the asrolled sheet. Similar to the as-rolled sheet in this work, there is only a small dependence of the mechanical properties in the Spectrulite and Russian material. The variation of the mechanical behaviour between the investigated directions is nearly the same for all sheets. Table 2. Tensile properties of AZ31 for the various conditions

material

direction

σ0.2 [MPa]

UTS [Mpa]

eu [%]

EL [%]

AZ31 as-extruded AZ31 as-rolled

longitudinal long transverse 0° 45° 90° 0° 45° 90° 0° 45° 90

210 140 200 195 196 198 191 180 186 174 169

280 255 250 252 257 269 267 265 263 259 259

n. a. n.a. 15.3 16.5 16.1 15.7 17.0 18.1 16.3 17.3 18.4

15.0 11.0 19.5 22.0 21.9 24.6 26.6 26.8 26.2 25.6 27.9

Spectrulite sheet Russian sheet

443 Corresponding to the very weak anisotropy of mechanical properties of the sheets the fracture surfaces after tensile tests (Figure 6) shows no significant differences between the three directions.

a) 0° (rolling direction)

b) 45°

c) 90° (transverse direction)

Figure 6. SEM pictures of fracture surfaces of the as-rolled AZ31 sheet

4

Conclusions

The anisotropy of the mechanical behaviour of AZ31 can be removed by rolling. All investigated sheets show an isotropic or near isotropic behaviour inside the rolling plane. This is corresponding to the measured pole figures and the picture of the fracture surfaces. The values for the stength of all sheets are nearly the same as for the extruded material. The values for elongation are much higher for the sheet material comparing to the extruded AZ31. There is a variation of these values between the three investigated sheets.

5

Acknowledgments

The authors woud like to thank Otto Fuchs Metallwerke Meinerzhagen, Germany and Honsel Profilprodukte Soest Gemany for providing the magnesium extrusions. The work was done with financial support of the BMBF in the frame of the projekt "Magnesium im Automobilbau (MIA)".

6

Literature

[1] H. Friedrich, S. Schumann, Materialwissenschaft und Werkstofftechnik, 32, 2001, p. 6. [2] H. Friedrich, S. Schumann, Magnesium 2000 (Ed.: E. Aghion, D. Eliezer) MRI, 2000, p. 9. [3] A. Luo, Magnesium Technology 2000 (Ed.: H. Kaplan, J. Hryn, B. Clow), TMS, Warrendale, USA, 2000, p.89. [4] T. K. Aune and H. Westengen, Magnesium Alloys and their Applications (B. L. Mordike and F. Hehmann, eds.) DGM (1992), p. 221. [5] E. F. Emley, Principles of Magnesium Technology, Pergamon Press, Oxford, 1996, p. 550. [6] C. S. Roberts, Magnesium and its Alloys, Wiley, New York, 1960, p. 171. [7] W. Bulian, E. Fahrenhorst, Metallographie des Magnesiums, Springer-Verlag, Berlin, 1942, p. 70.

444

High Cycle Fatigue of Magnesium Alloys Herwig Mayer1, Anton Stich2, Bernhard Zettl1 and Hans-Günther Haldenwanger2 1 2

Universität für Bodenkultur, Vienna, Audi AG, Ingolstadt

1

Introduction

Magnesium alloys have received increased attention as possible structural materials because of their low density and their rather high strength-to-weight ratio. Automotive applications are one of the possibilities where the benefits of weight reduction and, consequently, fuel saving are important. Fatigue properties of materials used for stressed components have to be known, since cyclic loading rather than static loads may cause failure. The number of load cycles in automotive applications may be relatively high, e.g. several 108 cycles in wheels, and the high cycle fatigue properties are therefore of great interest. Fatigue data of several magnesium alloys in the regime below 107 or 108 cycles are compiled in Ref. [1]. In this work, fatigue properties of aluminum and zinc containing magnesium alloys are studied in the regime of approximately 106 to 109 cycles to failure. Constant amplitude fatigue properties in ambient air are compared to the performance of the technically widely used age hardened aluminum alloy AlMgSi1. Fatigue data are obtained with the ultrasonic fatigue testing method at loading frequency 20 kHz, which allows measurements in the high and very high cycle regime within reasonable testing times. Fatigue properties as well as the mechanisms of crack initiation will be discussed in wrought, forged, cast and fiber reinforced AZ-magnesium alloys.

2

Materials and Methods

Fatigue experiments were performed with the wrought magnesium alloys AZ31hp (3%Al1%Zn-0,3%Mn) and AZ61hp (6,5%Al-1%Zn-0,3%Mn). The alloy AZ80hp (8,5%Al0,5%Zn-0,2%Mn) was tested in wrought and forged condition. The magnesium alloys were not heat treated after extrusion or forging. The aluminum alloy AlMgSi1 (AA6082) serves for comparison and to rank the magnesium alloys within the lightweight metals. Wrought AlMgSi1 was tested in age-hardened condition (T6). Additionally AlMgSi1 was tested as forged material, with solution treatment and artificial ageing (T6) after forging. The cast alloy AZ91 (9,5%Al-0,5%Zn-0,3%Mn) is used to investigate the influence of different casting techniques on the fatigue performance. In one series of experiments, specimens have been obtained from bars produced by high-pressure die-casting. This casting technique produces porosity (shrinkage and gas porosity), and fatigue cracks initiated at voids in all specimens (mean void diameter 0,59 mm, maximum void diameter 1,10 mm). Vacural die-casting is an appropriate casting technique to significantly reduce porosity and served to study the fatigue behavior of AZ91hp with limited deleterious

445 influence of voids. Additionally, squeeze cast AZ91 with 20 mass-% randomly oriented discontinuous aluminumoxide fibers (Saffil) has been tested as representative of magnesium matrix composites. The surfaces of all specimens were polished with abrasive paper of grade 1000 parallel to the specimens’ axis prior to the fatigue experiments. Fatigue data were measured with the ultrasonic fatigue testing method [2]. With this method, the specimens are excited to longitudinal resonance vibrations at cycling frequency of approximately 20 kHz, which leads to cyclic straining of the specimens and eventually crack initiation and failure. Strain gauges are used for calibration and monitoring the cyclic strain during the experiments, and stress amplitudes are calculated using the measured cyclic strain and the Young´s modulus, i.e. 65 GPa for fiber reinforced AZ91 and 45 GPa for the other magnesium alloys. Linear elastic behavior of the material is assumed, which is reasonable in the high and very high cycle fatigue regime where plastic strain amplitudes are small. Vibration amplitude of specimen’s end is used to control the magnitude of loading, i.e. the experiments are displacement controlled and the vibration amplitude is kept constant. Initiation of a fatigue crack increases the compliance of the specimen and reduces its resonance frequency. Control of resonance frequency is used for automatic operation and to detect 9 specimen failure. Specimens were cycled at least 10 cycles if they did not fail. Fatigue experiments were performed at constant amplitude and fully reversed tensioncompression loading (load ratio R = -1). Testing environment was ambient air of 20-22 °C and 40-60% relative humidity. Intermitted ultrasonic loading, i.e. sequences of pulses and pauses and forced air-cooling served to keep specimens´ temperatures below about 25 °C.

3

Results and Discussion

S-N fatigue data of wrought magnesium alloys AZ31hp and AZ61hp are compared with the cyclic properties of wrought aluminum alloy AlMgSi1-T6 in Figure 1. Specimens, which did not fail (runnouts) are marked with an arrow. Figure 2 shows the results of S-N tests with wrought AZ80hp, forged AZ80hp and with forged AlMgSi1-T6. Fatigue data of wrought AZ80hp have been obtained from defect free specimens. In approximately 35% of the actually tested wrought AZ80hp specimens, the crack initiated in regions containing numerous small voids (diameter 2-8 μm) created during extrusion, and these specimens are not considered in Figure 2. Static strength properties of the investigated materials and stress amplitudes at mean lifetimes of 107 and 5 x 108 cycles to failure are summarized in Table 1. In the regime of approximately 106-109 cycles the S-N data of the wrought and forged alloys can be reasonably well approximated using straight lines in double logarithmic plot, and no fatigue limit is found. Lines in the S-N plots indicate a fracture probability of 50%. The best fatigue properties among the tested wrought alloys shows (defect free) AZ80hp. In the regime above 107 cycles to failure this material shows cyclic properties, which are superior to wrought AlMgSi1-T6. Mean fatigue limit at 107 cycles of AZ61hp and AZ31hp is approximately 72% and 83%, respectively of the cyclic strength of wrought AlMgSi1-T6. Forging lowers the fatigue strength of AZ80hp to 75-80% of the 8 cyclic strength of wrought AZ80hp. Below approximately 10 cycles to failure, forged AZ80hp shows lower cyclic strength than forged AlMgSi1-T6, however comparable fatigue properties were found in the very high cycle regime. Considering the mass density of magnesium alloys, which is approximately 2/3 of aluminum alloys, cyclic strength divided by mass density of the investigated wrought and forged magnesium alloys is similar or higher than of AlMgSi1-T6.

446 Table 1. Tensile strength, yield stress and stress amplitudes at mean lifetimes of 10 and 5 x 10 cycles, respectively (all stresses in MPa). 7

Static Strength, Rm

Material

Yield Stress, RP0.2

8

50% Fracture 50% Fracture Probability at Probability at 107 Cycles 5 x 108 Cycles

Wrought AZ31hp

260

170

114

89

Wrought AZ61hp

300

220

132

101

Wrought AZ80hp

350

280

167

139

>310

>260

159

123

Forged AZ80hp

320

200

127

111

Forged AlMgSi1-T6

350

330

146

104

High-Pressure Die Cast AZ91hp

230

140

45

45

Vacural Die Cast AZ91hp

-

-

85

67

Squeeze Cast AZ91 + 20% Saffil

210

-

113

77

Wrought AlMgSi1-T6

Stress Amplitude (MPa)

200

150

120 100 Wrought AZ31hp Wrought AZ61hp

80 70

Wrought AlMgSi1-T6 105

106

107 108 Number of Cycles

109

Figure 1. S-N data of wrought AZ31hp (closed circles), wrought AZ61hp (closed squares) and wrought AlMgSi1-T6 (open circles).

447

Stress Amplitude (MPa)

200

150

120

100 Wrought AZ80hp Forged AZ80hp

80

Forged AlMgSi1-T6 70

105

106

107 108 Number of Cycles

109

Figure 2. S-N data of wrought AZ80hp (closed circles), forged AZ80hp (closed squares) and forged AlMgSi1-T6 (open circles).

Fatigue crack initiation in the wrought magnesium alloys is at the surface in most specimens (>95%) and internal crack initiation is rare. The crack initiates transgranular, and smooth and flat cleavage fracture is found in the first 50 to 100 μm, which mainly determine lifetimes in the very high cycle regime. Longer cracks propagate mainly transgranular and the fracture surfaces become more structured and rough. The forged alloy AZ80hp shows pronounced texture. Layers with approximate thickness of 100 μm are oriented about parallel to the direction of the applied principal stress, i.e. in length direction of the specimen. The interfaces between these layers are highly plastically deformed and may show small voids and cracks in some specimens. Interface regions are preferential sites for crack initiation, and cracks are formed at the surface of the specimen as well as (less frequent) in the interior. Fatigue data of high-pressure die cast AZ91hp obtained at 20 kHz are shown in Figure 3 with closed circles. This is compared to reversed tension compression fatigue data from the literature measured at 50 Hz using material of the same production (open circles) [3]. Within the ranges of scatter the results at both cycling frequencies are similar. This means that in the high cycle regime frequency effects (strain rate influences, corrosive influences of air humidity) and control of cyclic loading, i.e. force vs. displacement controlled experiments, are of minor influence on the obtained results, and high frequency data are applicable for components stressed at far lower frequencies in service. Additionally, fatigue data of Vacural die cast AZ91hp and of 20 % Saffil fiber reinforced AZ91 are shown in Figure 3. Data of Vacural die cast material refer to specimens with low porosity. Specimens were fatigue crack initiated at large porosity and where lifetimes were significantly smaller are not included.

448

120

Stress Amplitude (MPa)

100 80 70 60 50 40

30 104

High-Pressure Die Cast AZ91hp Vacural Die Cast AZ91hp AZ91 + 20 % Al2O3-Fiber 105

106 107 108 Number of Cycles

109

Figure 3. S-N data of high-pressure die cast AZ91hp determined at 20 kHz (closed circles) and 50 Hz (open circles) [3], of Vacural die cast AZ91hp (open squares) and of 20% aluminumoxide discontinuous fiber reinforced AZ91 (closed squares).

The mean fatigue strengths at 107 cycles of the investigated wrought and forged magnesium alloys are approximately 40 to 50% of the tensile strength, whereas it is approximately 20% for high-pressure die cast AZ91hp (Table 1). In this material fatigue cracks initiate at porosity in all specimens. Stress concentrations caused by voids accelerate crack initiation and fatigue lifetime is mainly determined by the cycles necessary to propagate the crack to rupture. Lifetimes of specimens and size of porosity at the crack initiation place may be correlated, and data scatter is mainly caused by the varying porosity in the different specimens. Below approximately 106 - 107 cycles to failure, the S-N curve may be reasonably well approximated by a straight line, whereas failures at higher number of cycles are rare. At low cyclic stresses fatigue cracks may initiate, however, stop propagating and do not cause rupture of the specimen [4]. The mean fatigue limit of high-pressure die cast AZ91hp is 45 MPa, which is approximately 60 % of the mean fatigue limit of similarly produced high-pressure die cast aluminum alloy AlSi9Cu3 [4]. In absence of voids, e.g. in Vacural die cast AZ91hp, fatigue cracks initiate at the surface. S-N data may be approximated with a straight line in a double logarithmic plot and no fatigue limit is found. Reinforcing AZ91 with 20% Saffil fibers increases the Young’s modulus by 44%. Comparing the stress amplitude at mean lifetimes of 107 cycles, the squeeze cast composite shows 33% higher fatigue strength than Vacural die cast AZ91hp. At 5 x 108 cycles, fiber reinforced AZ91 shows 15% higher fatigue strength. The increase of stiffness does not lead to similar increase of the fatigue strength at very high numbers of cycles.

449 One reason is the distribution of fibers in the magnesium matrix, where regions with low fiber density act as preferential crack initiation sites. Crack initiation in other specimens was found at bundles of fibers oriented normal to the applied principal stress where the interfaces of fibers and magnesium matrix failed. Regions where the infiltration of the fibers with AZ91 was poor also served as preferential crack initiation places. Crack initiation was found in the interior as well as at the surface of the specimens. The higher the number of load cycles the smaller material defects may influence the fatigue properties, which is considered as main reason for the relatively steep slope of the S-N curve of the composite in the very high cycle regime.

4

Conclusions

AZ-magnesium alloys produced as wrought, forged, cast and fiber reinforced materials have been tested at ultrasonic frequency at constant amplitude in the regime of approximately 106 - 109 cycles to failure. 1. Mean cyclic strength of wrought AZ31hp and AZ61hp is approximately 70% or 80%, respectively, of the cyclic strength of wrought aluminum alloy AlMgSi1-T6 and wrought AZ80hp shows similar or better fatigue properties. 8 2. Fatigue properties of forged AZ80hp at 10 cycles to failure or above are similar to forged AlMgSi1-T6, whereas the aluminum alloy shows superior fatigue properties at lower numbers of cycles. 3. The fatigue behavior of cast AZ91 is strongly affected by porosity. High-pressure die cast material showed a mean fatigue limit at 45 MPa whereas no endurance limit and superior fatigue properties were found in Vacural die cast material. 4. Squeeze cast Saffil fiber reinforced AZ91 showed better fatigue properties than not reinforced AZ91, more pronounced at lower number of cycles to failure.

5

References

[1] R.I. Stephens, C.D. Schrader in Magnesium and Magnesium Alloys (Ed.: M.M. Avedesian, H. Baker), ASM Speciality Handbook, Materials Park, OH, 1999, 249-257 and 264-275 [2] H. Mayer, Int. Mater. Rev. 1999, 44, 1, 1-34 [3] G. Schindelbacher, R. Rösch in Proc. Magnesium alloys and their applications (Ed.: B. L. Mordike and K. U. Kainer) Werkstoff-Informationsgesellschaft, Frankfurt, 1998 [4] U. Fuchs, H. Lipowsky, H. Mayer, M. Papakyriacou, A. Stich, S. Tschegg, B. Zettl, Mat.-wiss. u. Werkstofftech. 2002, 33, 3, 117-127

450

Fatigue Life Prediction of Magnesium Alloys for Structural Applications Yukio Miyashita, Zainuddin bin Sajuri, Takashi Umehara and Yoshiharu Mutoh Nagaoka University of Technology, Nagaoka, Japan

1

Introduction

Since high fatigue reliability of magnesium alloys is essential to their structural applications, it is relevant to make clear their fatigue characteristics. In the previous study on fatigue behavior of an extruded AZ61 under wide range of ambient humidity conditions [1], it was found that corrosion pits were formed and subsequently fatigue cracks were nucleated from the pits at applied stresses below the fatigue limit under higher humid condition. Die cast magnesium alloys have been widely utilized in many applications due to their advantages, such as near-net-shape forming, reduction of welding and machining processes and low cost. Inherent flows in die cast materials, size and shape of which depend on the casting process, significantly reduce the fatigue strength. In the present study, fatigue life prediction of magnesium alloys based on fracture mechanics approach has been made, where the pit depth for an extruded AZ61 at high humidity and the inherent defect diameter for a die cast AZ91D have been assumed as the initial crack size.

2

Experimental procedure

In order to study the effect of environment, fatigue strength test was carried out by using round bar specimens of extruded AZ61 under 20 °C -55%RH and 50 °C -80%RH. Effect of environment on tensile property was not so significant as shown in Table 1. Fatigue loading was applied with sinusoidal wave form with stress ratio, R = -1 and frequency of 20 Hz. A detail of a procedure of fatigue strength test is shown in the reference [1]. A center cracked plate tension (CCT) specimen was used for fatigue crack growth test. EDM notch was used as a pre-crack starter. Fatigue loading and environment conditions applied were the same as in the fatigue strength test. The crack length during fatigue crack growth test was measured by replica technique. A strain gage was mounted ahead of the crack tip on the specimen surface to evaluate crack closure behavior. Die cast and extruded AZ91D plates were used to study the effect of material processing. Mechanical properties of the materials used are shown in Table 2. Specimens were polished by using up to 1500 grit emery paper. Testing environment was 20 °C-55% RH. Cyclic load was applied under load control of sinusoidal wave form with stress ratio, R = 0.1 and frequency of 20 Hz. Fatigue crack growth test was carried out with center notched specimen. A pre crack was initiated from 1mm diameter hole at the center of the specimen by applying fatigue load. Crack closure behavior was evaluated by using strain gage attached ahead of the crack tip during fatigue crack growth tests.

451 Table 1. Tensile properties of extruded AZ61.

Temperature & Humidity

Young’s modu- Tensile strength lus σb [MPa] E [GPa]

20 °C-55%RH 43 50 °C-80%RH 44

329 326

0.2% Proof stress Elongation σ0.2% [MPa] [%] 244 240

23 28

Table 2. Tensile properties of die cast and extruded AZ91D at room temperature.

Young’s modulus Tensile strength 0.2% Proof stress Elongation σ0.2 [MPa] [%] E [GPa] σb [MPa] Die cast Extruded

45 41

230 320

160 195

3 18

3

Results and discussions

3.1

Fatigue life prediction for an extruded material with corrosion pits

From the fatigue test results [1] shown in Fig.1, it was found that fatigue strength was considerably reduced at applied stresses below the fatigue limit and fatigue failure occurred from pits. Therefore, it is suggested that, the conjoint action between localized plastic deformation due to cyclic load and environment leads to formation of pit and subsequent fatigue crack initiation under high humidity condition at stresses below the fatigue limit. Existence of corrosion pit at the fracture origin could be observed for all the specimens tested at stress amplitudes lower than the fatigue limit under 80%RH. An example of corrosion pit observed at the fracture origin is shown in Fig.2. These observations confirm that fatigue cracks initiate from the corrosion pits at stress levels lower than the fatigue limit at high humidity condition. Formation of corrosion pit and crack initiation behavior was observed by using replica technique during the fatigue test. Pit and crack initiation lives are shown in Fig.1 at applied stresses of 130 and 140 MPa under 50 °C -80%RH. From the results, the crack initiation life was about 50% of total fatigue life. 20͠-55㧑RH 50͠-80㧑RH Interrupted test (50 ͠-80㧑RH) Pit initiation life Crack initiation life Fatigue life

Stress amplitude, ǻa (MPa)

200 180

(a)

160 140 120 100

P

μO P

P

P

P

P P

(b)

P

103 104 105 106 107 Number of cycle to failure, N f (cycles)

Figure 1. Result of fatigue test of AZ61 at 20 °C55% RH and 50 °C-80% RH.

μO

Figure 2. A corrosion pit as a crack origin of the specimen fractured at 130 MPa. (50 °C -80% RH). (a) fracture surface, (b) specimen surface.

452 Figure 3 shows the effect of environment on fatigue crack growth behavior of extruded AZ61 alloy. From the results, it is found that the fatigue crack growth curves under 20 °C -55% RH and 50 °C -80% RH were almost identical. Therefore, it is suggested that the crack growth behavior is not significantly influenced by ambient humidity. In this study, the threshold values for ΔK and ΔKeff are 1.34 and 0.4 MPam1/2, respectively. The da/dNΔK curves obtained in the Paris regime can be expressed as: da m (1) = C ( ΔK ) dN where C is a material constant and m is the slope of the curve. C and m values for ΔK and ΔKeff crack growth curves are C = 1.1 × 10-9m/cycle, m = 2.0 and C = 7.6 × 10-9 m/cycle, m = 1.5, respectively. Total fatigue life prediction at 50 °C -80%RH was obtained by the sum of predicted pit growth life and fatigue crack growth life based on Paris’s law. Kondo [2] has suggested that the pit size has almost remained the same size after the transition from pit to crack. This phenomenon occurred probably because the anodic site moved from the pit. According to this assumption, the pit depth observed at the fracture origin can be considered as the critical pit depth at which the transition from pit into crack occurred. In fracture mechanics approach to fatigue crack growth life prediction, the critical pit depth is assumed as the initial crack size [3]-[7]. Considering the effect of environment, the corrosion pit growth law is given as [8]; 1

⎛ N ⎞3 a = Cpt = Cp ⎜ ⎟ (2) ⎝ f ⎠ Here, a is the pit depth, t is the time, N is the number of loading cycles, f is the frequency and Cp is the coefficient. As a fatigue crack initiated from a corrosion pit at almost half of the total fatigue life, the time spent for pit growth to a critical pit depth would be also nearly half of the total fatigue life. The diameter 2c and depth of the pit b on the fracture surface were also measured. The aspect ratio b/c was about 0.9. By considering corrosion pits as surface cracks, the stress intensity factor at root of the pit can be calculated by using the equation proposed by Newman and Raju [9]: 1 3

(ΔK ) pit = 1.1σ a π b

Q where σa is the stress amplitude, a is the pit depth and Q is the shape factor. The pit growth rate is given as;

(3)

3

Cp Cp ⎛ da ⎞ (4) ⎜ dN ⎟ = 13 23 = 3 fa 2 ⎝ ⎠ pit 3 f N By substituting Equation (3) into Equation (4), the relationship between (da/dN)pit and (ΔK)pit can be expressed as;

C p π 2 (1.1σ a ) ⎛ db ⎞ = ⎜ dN ⎟ 4 ⎝ ⎠ pit 3 fQ 2 ( ΔK ) pit 3

4

(5)

By using these equations, pit growth rate curves db/dN-ΔKeff at stress amplitudes of 110, 130 and 150 MPa were plotted using solid lines in Fig. 4. From the figure, as the pit

453 grows, the pit growth rate decreases and the fatigue crack growth rate increases. Finally, fatigue crack growth rate exceeds the pit growth rate. The intersection points between these two curves are the transition points from pit growth into crack growth. Moreover, the higher the stress amplitude, the higher the stress intensity factor at fatigue crack nucleation point. Fatigue crack growth life after crack nucleation can be predicted based on the linear relationship between crack growth rate and effective stress intensity factor in Region II, as follows: 2 1− m 1− m (6) N crack = b f 2 − bcr 2 m ⎛ ⎞ π ( 2 − m ) C ⎜ 1.1σ a Q ⎟ ⎝ ⎠ Here, bcr and bf are the critical pit depth and the fatigue crack depth at fracture, respectively. The fatigue crack depth at fracture bf is defined as the depth of semi-elliptical crack at which the stress intensity factor reaches the critical value. The summation of predicted pit growth life and crack growth life will be the total fatigue life under test environment of 50 °C -80 % RH. The predicted lives for pit growth and the fatigue crack growth are schematically shown in Fig. 5 with a dotted and solid line, respectively. From the figure, at low stress amplitude levels, the predicted lives agree with the experimental results. However, at high stress levels near 150 MPa, the predictions overestimate the lives. The small crack effect may be one of the reasons for the overestimation of lives.

(

10-6

ǍK

-7

10-6

20͠-55㧑RH 50͠-80㧑RH

ǍKeff 20͠-55㧑RH 50͠-80㧑RH

Corrosion pit growth rate, da/dN [m/cycle] Fatigue crack growth rate, da/dN [m/cycle]

Crack growth rate, da/dN [m/cycle]

10

)

10-8 10-9 10-10 10-11

10-7 10-8 10-9 10-10 150MPa 130MPa 110MPa

10-11 10-12

10-12 10-13

10-13 0.1 0.2 0.4

1

2

4

10

Fatigue crack Corrosion pit

0.1 0.2 0.4

1

2

4

10

20 1/2

Effective stress intensity factor range, ٌKeff [MPa.m ]

20 1/2

Stress intensity factor range, ٌK [MPa.m ]

Figure 3. Fatigue crack growth curves for AZ61 at 20 °C-55% RH and 50 °C-80% RH.

Stress amplitude, ǻa [MPa]

200

Figure 4. Corrosion pit and fatigue crack growth rate verses stress intensity factor at 50 °C-80% RH.

20͠55㧑RH 50͠80㧑RH Pit growth life (Cal.) Total fatigue life (Cal.)

180 160 140

P

P

P P

P

120 100 103

104 105 106 107 Number of cycles to failure, Nf [cycles]

108

Figure 5. Prediction of fatigue life for extruded AZ61 at 50 °C 80% RH.

454 3.2

Fatigue life prediction for die cast material

Results of fatigue strength tests of die cast and extruded AZ91D magnesium alloys are shown in Fig. 6. Fatigue life of the die cast material was shorter than that of the extruded material. An example of fracture surface of the die cast material is shown in Fig. 7. Fatigue crack initiated from an inherent flaw in all specimens of the die cast material, but not for the extruded material. Fatigue crack curve of the extruded material is shown in Fig. 8. ˂Kth = 1.2 MPam1/2 and parameters of crack growth curve according to Paris’s power law (Eq. (1)) C = 5.52 × 10-10 and m = 3.14 are obtained from the figure. Fatigue life was predicted based on the fatigue crack growth curve, assuming that fatigue life is equal to crack growth life from the initial crack size to the critical size at final failure. A square root of projective area of the flaw in the principle stress direction was assumed as an initial crack size, i.e. area1/2 [10]. From the fracture surface observations, the smallest, largest and average area1/2 obtained were a0min = 50, a0max = 250 and a0ave = 140 ˩m, respectively. It is also assumed that final failure occurs when ˂K reaches to 7.0MPam1/2. Predicted life is shown in Fig. 9. It shows good agreement between experimental and predicted lives. The predicted lives may be improved if the small crack effect is taken into consideration. Stress amplitude, ǻa [MPa]

100 80

ǴO

60 40 20

Extruded Die cast

0 3 10

104

105

106

107

108

ǴO

Number of cycles to failure, N f [cycles]

Figure 6. S-N curves for die cast and extruded AZ91D.

Figure 7. Fatigue fracture surface of the die-cast AZ91D (ǻa = 42.8MPa, Nf = 6003270 cycles). 70

10

Stress amplitude, σa [MPa]

Crack growth rate, da/dN[m/cycles]

10-6 ΔK ΔKeff

-7

10-8 10-9 10-10 10-11

50 40 30 103

10-12 10-13

60

0.1

0.2

0.4

1

2

4

Experimental results Die-cast Predicted life a0 min =50μm a0 max =250μm a0 ave =140μm

104 105 106 107 Number of cycles to failure, Nf [cycle]

108

10

Stress intensity factor [MPa.m 1/2]

Figure 8. Fatigue crack growth curves for extruded AZ91D.

Figure 9. Prediction of fatigue life for die cast AZ91D.

455

4

Conclusions

Fatigue lives for two different cases were successfully predicted. One is for the case where a fatigue crack nucleates from a corrosion pit. The other is for the case where a fatigue crack initiates from a inherent flaw. Summaries of the present work are as follows: 1) At the corrosion fatigue region where high relative humidity, fatigue fracture was originated from corrosion pits. The corrosion pit grew up to a critical size for initiating a fatigue crack. 2) Fatigue crack growth curves for AZ61 under both 20 °C-55% RH and 50 °C-80% RH test environments were almost identical. Therefore, it is suggested that the crack growth behavior is not significantly influenced by ambient humidity. 3) The predicted lives agree with the experimental results. However, at high stress levels near 150 MPa, the predictions overestimate the lives. The small crack effect may be one of the reasons for the overestimation of lives. 4) Die cast AZ91D shows lower fatigue strength compared to extruded AZ91D. This is dominantly due to existence of the inherent flaw in die cast material that serves as a fatigue crack starts. 5) Predicted fatigue lives of the die cast material was good agreement with experimental results, where the total fatigue lives was assumed to be almost equal to the fatigue crack growth life started from a inherent flow

Acknowledgement The present research work was partly supported by New Energy and Industrial Technology Development Organization (NEDO) under the Ministry of Economy, Trade and Industry (METI), as well as by 21st Century COE program of the Ministry of Education, Culture, Sports, Science and Technology of Japan.

References [1] Z.B. Sajuri, Y. Miyashita and Y. Mutoh, J. Japan Inst. Light Metals, 52 (4), 161 (2002). [2] Y. Kondo, Corrosion, 45 (1), 7 (1989). [3] K.K. Sankaran, R. Perez and K.V. Jata: Mater. Sci. and Eng. A297, 223 (2001). [4] G.S. Chen, K.C.Wan, M. Gao, R.P. Wei and T.H. Flournoy: Mater. Sci. and Eng., A219, 126 (1996). [5] D. Kujawski, Int. J. Fatigue, 23, 95 (2001). [6] D.W. Hoeppner, Fatigue Mechanisms, ASTM STP 675, ASTM, Philadelphia, 841 (1979). [7] S.I. Rokhlin, J.Y. Kim, H. Nagy, B. Zoofan, Eng. Fracture Mechanics 62, 425 (1999). [8] Hugh P. Godard, The Canadian Journal of Chemical Engineering, 167 (October 1960). [9] I.S. Raju and J.C. Newman, ASTM Special Technical Publication, 789 (1986). [10] Y. Murakami, T. Toriyama, Y. Koyasu and S. Nishida, TETSU-TO-HAGANE, 79-6, 678 (1993).

456

Correlation of Microstructure and Mechanical Properties of Rolled Magnesium Sheet AZ31 F. Kaiser, J. Bohlen, D. Letzig, K.U. Kainer1) A. Styczynski, Ch. Hartig2) 1) 2)

GKSS Research Center GmbH, Geesthacht Technical University Hamburg-Harburg, Hamburg

1

Introduction

The main reason for applying light-weight metals in the automotive industry is the potential of saving natural energy resources simply by weight reduction. Magnesium as the lightest available construction metal offers a wide range of opportunities for the use in automobiles. However, presently mostly cast parts are applied the use of wrought magnesium alloys is still in an underestimated position. This is mainly due to two aspects, first there is the lack of sufficient cold formability and anisotropy of the final sheets. The rolling process itself on the other hand is not adequately developed either and still needs a lot of improvement in order to produce sheets with properties demanded by the automobile manufacturers. Generally, the mechanical behaviour is affected by changes in the grain size, the microstructural homogeneity and the texture. Focusing on a plastic deformation process like rolling, the process parameters cause these changes in the microstructure which again will result in decisive mechanical properties. The complex mechanisms of deformation, recovery and recrystallisation during deformation at elevated temperatures are important for understanding the needs of a process development. The present paper prosecutes two aspects of the deformation behaviour of magnesium sheets. First it will give a brief mechanical characterization of a commercial magnesium with a distinct focus on the anisotropy and the possible microstructural causes of this behaviour. The rolling process with its influence on the sheets in the first place will be part of the second aspect. The aim is to show the influence of the parameters on the microstructure and texture and how this could affect the anisotropy in the rolled sheets.

2

Experimental Procedure

A commercial AZ31B-rolled sheet (with nominal composition 3 wt% Al, 1 wt% Zn, balance Mg) was provided in the stress-relieved H24 temper and thicknesses of 1.0 and 1.6 mm. The tensile tests at room temperature were conducted on a ZWICK® Z050 universal testing machine with a constant strain rate of 10-3 s-1. For the elevated temperature test, a standard forced-air furnace was attached to the machine. The tensile samples had a reduced section sized 125 x 20 x 1.6 mm, respectively 125 x 20 x 1.0 mm. Texture measurements have been taken out using x-ray diffraction in reflection geometry. The metallography was performed by picric acid etching.

457 A laboratory mill was used for all rolling experiments. The rolls were kept at room temperature and operated with a speed of 3 m/min as rolling rate. The specimens for the stepwise rolling had a size of roughly 50 x 50 x 10 mm (squeeze cast) and 100 x 50 x 10 mm (gravity casting) respectively. For the one-step rolling the specimens hat all different geometries in order to reach a final sheet thickness of 3mm each. All rolling was done with a specimen temperature of 400 °C at the time of withdrawal from the furnace, which was located directly next to the mill. The first part of the experimental processing consisted of a stepwise rolling by reducing the thickness in logarithmic steps of 0.1. For this experiment two different feedstock samples of AZ31, a squeeze-cast specimen and a specimen produced by permanent mould casting were compared. After every 5 steps a sample was taken from the rolled material after air-cooling and kept for further analysis. The remaining material was annealed for 20 min before further rolling in order to control the material temperature. It also was treated 2 min between two rolling passes. Thus a rolling process results which leads to thin sheets up to a final reduction in thickness of at least 86% (ϕ = 2). For the second part of the experiments AZ31 slab cast specimens with were also rolled at 400 °C, but this time in one step directly down to a final sheet thickness of 3 mm, which equals deforming degrees of ϕ = 0.17 (16%) to 1.05 (65%). The material was solution annealed for 10 hours before rolling to avoid any influence from precipitations and water-quenched after rolling.

3

Experimental Results

3.1

RT Tensile Tests

3.1.1 As-received condition First measurements were tensile tests at room temperature in order to obtain the mechanical properties such as yield strength and tensile strength a function of the sheet orientation For these tests, specimens were taken from the sheet in 0° and 90° whereby 0° means rolling direction (RD) and 90° indicates transverse direction (TD). Fig. 1 clearly shows the high dependence of the YS and UTS on the orientation of the sheet. The yielding difference between rolling direction (0°) and transverse direction (90°) is about 50 MPa, which is roughly 20% of the maximum value at 90°. On the contrary the ultimate tensile strength shows a smaller difference. There seems to be no great effect of the anisotropy on ductility though the transverse direction shows a slightly higher strain in average. 3.1.2 Heat treated samples For the heat treatments, 1.6 mm RD- and TD-specimen were used and exposed to the respective temperatures for 30 minutes in a standard forced-air furnace. Afterwards, tensile tests were carried out at room temperature. While at room temperature, the difference of the YS between RD and TD is close to 60 MPa it drops down to roughly 30 Mpa between 150 and 200 °C heat treatment temperature (Fig.2). This difference remains almost constant throughout the higher temperatures. The anisotropy of the tensile strength follows the same pattern, only the differences are smaller from the beginning on and can be neglected with increasing heat treatment temperature.

458 350

80

a)

70

250

60

Transverse Direction

200

ΔYS [MPa]

Engineering Stress [MPa]

300

Rolling Direction

150 100

30

10 0

5

10

15

20

0

25

Engineering Strain [%]

Fig. 1: Stress-strain-curves for rolling direction and transverse direction, tests at RT and ε& =10-3 s-1, 1.0mm sheet

3.2

40

20

50 0

50

0 20 50 100 150 200 250 300 350 400 450 500 550

As-Received / Heat Treatment Temperature [°C]

Fig. 2: Yielding anisotropy ΔYS as a function of the testing temperature

Texture Analysis

3.2.1 Commercial Sheet The textures were measured for the 1.6mm sheet in as-received condition. Fig. 3 shows the (0002)-basal plane pole figure and the (10 1 0)-prismatic plane pole figure for the feedstock sheet, whereas Fig. 4 represents the corresponding pole figures for the 250 °C heat treated condition. The basal plane pole figures show a distinctive preferred orientation of the (0002)-plane for both conditions, the prismatic planes show no preferred orientation for either conditions. But it becomes obvious, that the (0002)-basal plane distribution orientation intensity of the heat treated sheet drops roughly 25% and is more centered compared to the original state.

Fig. 3: Pole figures of the (0002)- and (10 1 0)-planes (asreceived condition)

Fig. 4: Pole figures of the (0002)- and (10 1 0)planes (250 °C)

3.2.2 Rolled samples

Due to the nature of cast materials, the texture of the feedstock shows random orientation and is therefore not presented here. The basal plane and prismatic plane textures are given for all three materials for the deformation degree of roughly ϕ = 1.

459 It shows for both stepwise rolled castings that the preferred orientation of the basal planes basically don’t show an offset while the one-step rolled sample shows the typical double peak even though with unequal intensity. The (10 1 0)-prismatic plane distribution shows no significant conspicuity in comparison (Fig. 5, 6 and 7).

Fig. 5: Pole figures for squeeze cast sample (400 °C, ϕ = 0,96, multiple step rolling)

Fig. 6: Pole figures for gravity cast sample (400 °C, ϕ = 0,98, multiple step rolling)

Fig. 7: Pole figures for slab cast sample (400 °C, ϕ = 1,05, single step rolling)

3.3

Microstructure evaluation

Optical microstructure investigations were carried out on the 1.6 mm heat treated sheets and the rolled samples to characterize the evolution of the grain size with rising temperature and increasing degree of deformation. The grain size distribution of the commercial sheet and its heat treated condition is very inhomogeneous. The feedstock consists of grains full of twins, visible grains are 20 μm in average size, single grains go as high as 60 μm. The heat-treated specimen is free of twins, the size is approx. 5 μm in average. The decrease in the size after the heat treatment is probably due to the fact, that the microstructure of the as-received condition is hardly identifiable. The squeeze cast feedstock has an average grain size of roughly 300 μm, the gravity cast material shows a grain size of 1800 μm on average. The rolled sheets after a deformation of ϕ ≈ 1 have grains sized between 10 and 15 μm which is comparable to the commercial sheet. While with the slab cast feedstock sample an average grain size of 350 μm could be observed, the rolled samples don’t show any grain structure anymore. Instead one can see highly deformed grains which form some typical bands throughout the specimen oriented in rolling direction. In some areas show the starting of dynamic recrystallization (DRX). The new grains have sizes below 1 μm (Fig. 9).

460 a)

c)

e)

b)

d)

f)

Fig. 8: Commercial sheet in as-received condition (a) and heat treated @250 °C, 30 min (b) Squeeze casting and the gravity casting (c,e) and the corresponding materials in rolled condition (d,e), rolling stepwise @400 °C

Fig. 9: Micrographs of the slab cast feedstock (left), rolled condition (middle) and rolled condition with DRX (right); material rolled in a single step @400 °C

4

Discussion

The anisotropic behaviour of the AZ31-sheet at room temperature i.e. the differences in the stress-strain curves of RD and TD, in non-heat treated condition is mainly a result of the texture. For a first basic approach it can be assumed, that the slightly "egg-shaped" distribution of the (0002)-basal plane has a significant influence on the planar anisotropy of the sheet. We can understand the behaviour in a way, that there is a better chance for the activation of slip systems according to the Schmid factor in rolling direction than in the transverse direction. That might result in the activation in other slip systems. This would result in higher stresses necessary for the 90°-orientation, hence increasing the yield strength. The heat treated tensile-specimen show lower mechanical strengths, though they remain almost constant beyond a heat treatment temperature of 200 °C. The recovery of the microstructure also seems to have a positive effect on the activation of deformation mechanism which results in the decreasing anisotropy of the yielding. The change from the typical off-set-peak texture to a centered distribution of the heat treated sample must have a

461 distinctive influence on the anisotropy when applying the explanation from above. A majority of the basal planes which used to be located approx. 20° towards the rolling direction are now moved in alignment with the rolling plane. Since these planes parallel to the rolling direction cannot contribute to gliding, the prior advantage of the slip activation in rolling direction gets smaller, thus anisotropy decreases as well. Since the double peak is doubtlessly a result of the rolling process it needs to be examined what the reasons are and what processing parameters influence the formation of this deformation texture. The experiments with both stepwise rolled cast samples result in an almost concentric distribution of the basal planes whereas the single-step rolling shows again the typical double-peak texture, even though there is no strong “egg-shaped” alignment in rolling direction. As a consequence the interim annealings must be the reason for the change in texture. Looking at the micrographs it becomes clear that both stepwise rolled specimens show a recrystallized microstructure with remains of plastic deformation like twins. The third rolling specimen which was deformed in a single step has no visible grain structure anymore which means that there has been no static recrystallization. However, rudiment dynamic recrystallization could be observed but the amount of new grains is very small and an effect on mechanical behaviour is not likely. Combining the results of the microstructures of rolled and annealed sheets one could assume, that the double peak is mainly a deformation texture with a strong influence on the anisotropic behaviour whereas annealing causes a change from a double texture towards a centered basal plane distribution with decreasing anisotropy as a consequence thereof.

5

Summary

Tensile tests were carried out on a commercial AZ31 sheet in 1.0 and 1.6 mm thickness in order to investigate the anisotropic behaviour of the mechanical properties, especially the yielding. Tests at RT show a strong vertical and planar anisotropy of the sheet, which was explained with the strong rolling texture. Experiments with heat-treated specimen lead to a decrease of the yielding anisotropy by approx. half of its corresponding value at RT. Rolling experiments in a single step and stepwise with annealing between the steps were conducted at 400 °C. The results indicate that annealing between the rolling steps leads to a concentric basal plane distribution whereas a single step rolled specimen shows an almost typical double peak texture. This allows the assumption that the double peak texture is a pure deformation texture and results in a stronger anisotropic behaviour than rolled and annealed material. Further investigations need to give more information on the mechanisms of anisotropy and their relationship to texture and microstructure in order to find proper processes to avoid this behaviour in terms of sheet metal forming.

6

References

[1] E.F. Emley, Principles of Magnesium Technology, Pergamon 1966 [2] A. Beck: Magnesium und seiner Legierungen, Springer 1939

462 [3] S.R. Agnew in: H.I. Kaplan Ed. „Magnesium Technology“, TMS Annual Meeting Seattle 2002, p. 169 - 174 [4] K.U. Kainer (Ed.): Magnesium Alloys and their Applications, Wiley-VCH 2000 [5] J. Bohlen, D. Letzig, H.G. Brokmeier, A. Styczynski, Ch. Hartig, K.U. Kainer in: H.I. Kaplan Ed. „Magnesium Technology“, TMS Annual Meeting Seattle 2002, p. 165 – 167 [6] T. Mukai, H. Watanabe, K. Ishikawa, K. Higashi in: H.I. Kaplan Ed. „Magnesium Technology“, TMS Annual Meeting Seattle 2002, p. 137 – 140 [7] K. Lange (Ed.): Umformtechnik Band 3 – Blechbearbeitung, Springer 1990 [8] J. Bohlen, A. Horstmann, F. Kaiser, A. Styczynski, D. Letzig, K.U. Kainer in: H.I. Kaplan Ed. „Magnesium Technology“, TMS Annual Meeting Seattle 2003, p. 253 – 258 [8] V. Kree, J. Bohlen, D. Letzig, K.U. Kainer, Praktische Metallogrphie, 2003, in press

463

Effect of Warm Forming on Mechanical Properties of an AZ31 Magnesium Alloy Sheet H. Takuda, H. Fujimoto, D. Inoue Kyoto University, Kyoto

1

Introduction

Magnesium is poor in formability at room temperature due to the limited slip planes. However, with increasing temperature, the critical shear stress of nonbasal slip decreases and the number of slip planes increases [1]. AZ31 is one of the magnesium-based alloys that are commercially produced and most commonly used. It is reported that AZ31 sheets have good stretchability and drawability at elevated temperatures between 150 and 250 °C [2, 3]. Therefore, the warm forming of AZ31 sheets is considered to be practical use in recent years. However, few studies have examined the mechanical properties of the sheets after warm forming. In order to find the proper conditions of warm forming, the examination from such a viewpoint is indispensable. In this study, the influence of the temperature and the amount of deformation at warm forming on the tensile properties of an AZ31 sheet at room temperature after warm forming is investigated. The microstructural changes by warm forming are observed, and the relation to the tensile properties is examined. Furthermore, the tensile properties of the sheet during warm forming are evaluated.

2

Experimental Procedure

2.1

Material

Material used in this study is a commercial magnesium-based alloy AZ31B-O sheet with a thickness of 0.8 mm. Figure 1 shows the microstructure of the sheet. The grains are equiaxial and the average grain diameter is 11 μm. Rolling direction

20μm Figure 1: Microstructure of the sheet.

464 2.2

Tension Tests at Room and Elevated Temperatures

The tensile specimens were cut from the sheet in such a way as to be parallel (0°) and 90° to the rolling direction. The gauge length and width of the specimens were 50 and 12.5 mm, respectively. The specimens were previously elongated at 150, 200 and 250 °C simulating warm forming. Three kinds of elongations of about 4, 8 and 12% were given. And then, the tension tests were carried out for the specimens at room temperature. The cross-head velocity was constant at 3 mm min-1 for the above tests, corresponding to the initial strain rate of 8.3 × 10-4 s-1. In addition to the above tests the tension tests were carried out at room and elevated temperatures of 150, 200 and 250 °C, using the specimens with no previous deformation. The tests were carried out for three kinds of cross-head velocities of 3, 30 and 300 mm min-1 to examine the influence of strain rate.

3

Results and Discussion

3.1

Tensile Properties after Warm Forming

Figure 2 shows the true stress-strain curves obtained from the tension tests at room temperature after warm forming, compared to that without warm forming. The solid lines are the curves for the specimens in the direction of 0° to rolling, and the dotted ones for 90°. The horizontal axis of this figure indicates the total logarithmic strain including the strain at warm forming. The tensile properties obtained from these data are described below. Figures 3, 4 and 5 show the relationships between the amount of pre-forming at elevated temperatures and the proof stress, the tensile strength and the fracture elongation at room temperature, respectively. The tensile strength was evaluated using the cross section after warm forming as an initial one. For the warm forming at 250 °C, the change in proof stress, tensile strength and fracture elongation is very small from the properties before warm forming. As the temperature decreases, the influence of warm forming on the tensile properties becomes large. For the warm forming at 200 and 150 °C, the proof stress and the tensile strength increase and the fracture elongation decreases with increasing the amount of warm forming. The change in proof stress is large in particular. In case where the specimens are pre-deformed at 150 °C to the rolling direction with the elongation of 12%, the proof stress increases from 181 to 242 MPa, the tensile strength from 265 to 279 MPa, and the fracture elongation decreases from 25.2 to 19.3%. However, the total elongation including that at warm forming increases even for the case of 150 °C, as indicated in Figure 6. There is no qualitative difference depending on the tensile direction. However, it is obvious from the above figures that the stress is higher and the elongation is lower for the direction of 90°. The microstructures of the sheet after warm forming are shown in Figure 7. Compared to the microstructure of the sheet before deformation (Fig. 1), great changes are not observed in that deformed at 150 °C, namely the grain boundary is still obvious and linear. On the contrary, with increase in temperature, the grain boundary becomes complex line like serration, and in the microstructure deformed at 250° C small and recrystallized

465 grains are mingled among large grains. The above-mentioned result that the change in the tensile properties is very small for the warm forming at 250 °C is considered to be due to the partial recrystallization.

True stress, σ /MPa

300

200

100

0

0q 90q 0

0.1

0.2

0q 90q

0.3 0

T rue strain,ε

0.1

0.2

0q 90q

0.3 0

0.1

T rue strain,ε

(a) 150͠

0.2

0.3

T rue strain,ε

(b) 200͠

(c) 250͠

240

150 °C

220 200 250 °C

180 160

0 5 10 Elongation in warm forming / %

Elongation after warm forming / %

Figure 3: Relationships between proof stress after warm forming and elongation given at warm forming for various temperatures.

26

200 °C

24

250 °C

22 20 18

Tensile strength, σ /MPa

260

0° 90 ° 200 °C

0° 90°

150 °C

0 5 10 Elongation in warm forming / %

Figure 5: Relationships between fracture elongation after warm forming and elongation given at warm forming.

320 300

0° 90°

150 °C

280 260 240 220 200

200 °C

250 °C

0 5 10 Elongation in warm forming / %

Figure 4: Relationships between tensile strength after warm forming and elongation given at warm forming.

/%

280

Total elongation

Proof stress, σ0.2 /MPa

Figure 2: True stress-strain relationships at room temperature before and after warm forming at 150, 200 and 250 °C. The solid and dotted lines indicate values measured at tension tests in directions of 0 and 90° to rolling, respectively.

32 30

0° 90°

28 26 24 22

0 5 10 Elongation in warm forming / %

Figure 6: Relationships between total elongation and elongation given at warm forming for 150 °C.

466 20μm

Tensile direction

(a) 150 °C

(b) 200 °C

(c) 250 °C

Figure 7: Microstructures after warm forming at 150, 200 and 250 °C for an elongation of 12%.

3.2

Tensile Properties at Elevated Temperatures

The tensile properties of the sheet obtained from the tension tests at elevated temperatures are described below. Figure 8 shows the true stress-strain curves for various initial strain rates at room and elevated temperatures of 150, 200 and 250 °C. Note that our present curves deviate from the true curves after the occurrence of necking, because the stress and strain were evaluated under the assumption of uniform deformation. While there is little influence of the strain rate at room temperature, the stress and the elongation depend greatly upon the strain rate at elevated temperatures. It is recognized from the curves that necking occurs at elevated temperatures before fracture. The elongation increases with increasing temperature and decreasing strain rate. It is found that the sheet has sufficiently high ductility at the temperatures over 150 °C. 8.3˜10-2s-1

25͠

0q 90q

300

8.3˜10-2s-1 -3 -1 8.3˜10 s

-4 -1

8.3˜10 s

8.3˜10-2s-1 -3 -1 8.3˜10 s

200

200͠

100 -4 -1

0

(a) 0

8.3˜10 s

0.1 0.2 0.3 0.4 0.5 True strain, ε

True stress, σ /MPa

True stress, σ /MPa

300

200

150͠

0q 90q -4 -1

8.3˜10 s

-2 -1

8.3˜10 s

8.3˜10-3s-1

100

0

(b) 0

8.3˜10-4s-1

250͠

0.1 0.2 0.3 0.4 0.5 0.6 0.7 True strain, ε

Figure 8: True stress-strain relationships at room temperature (a) and elevated temperatures of 150 (b), 200 (a) and 250 °C (b) for various initial strain rates.

467

0.15

0q 90q

0.10 0.05 0

0

50

100 150 200 Temperature /͠

250

Normal anisotropy parameter, r

Figure 9: Relationships between strain rate sensitivity exponent, m, at a strain of 0.1 and temperature.

5 0q 90q

4 3

45q

2 1 0

0

50

100 150 200 Temperature /͠

250

Figure 11: Relationships between normal anisotropy parameter, r, and temperature.

Work-hardening exponent, n

Strain rate sensitivity exponent, m

Figure 9 shows the relationship between the temperature and the strain rate sensitivity exponent, m, evaluated at the strain of 0.1. The sensitivity increases with temperature and amounts to 0.15 at 250 °C. Figure 10 shows the relationship between the temperature and the work-hardening exponent, n, evaluated in the strain range between 0.03 and 0.15. The n-value decreases with increasing temperature. At elevated temperatures the n-value decreases with strain rate. The n-values for the tensile direction of 0° are somewhat larger than those for 90°. Figure 11 shows the relationship between the temperature and the normal anisotropy parameter, r, measured at the tension tests for the initial strain rate of 8.3 × 10-4 s-1. The r-value is very high at room temperature, particularly in the direction of 90°. However, the r-value decreases with increasing temperature and comes to be 1. 0.20

0q 90q

0.15 0.10

-4 -1

: 8.3˜10 s -3 -1 : 8.3˜10 s -2 -1 : 8.3˜10 s

0.05 0

0

50

100 150 200 Temperature /͠

250

Figure 10: Relationships between work hardening exponent, n, and temperature for various initial strain rates.

468

4

Conclusions

In this study, the influence of warm forming on the mechanical properties of the AZ31 magnesium alloy sheet has been investigated. Specimens were elongated in advance at 150, 200 and 250 °C, and tension tests were carried out at room temperature on the specimens. It has been clarified from microstructural observations that recrystallization occurs in the specimens deformed at 250 °C. The properties, such as proof stress, tensile strength and fracture elongation, of the specimens deformed at 250 °C are unchanged and are almost the same as those before warm forming. With decreasing temperature and with increasing elongation during warm forming, the proof stress and the tensile strength increase while the fracture elongation decreases. However, the decrease in ductility is comparatively small even for the specimens deformed at 150 °C. Tension tests were carried out also at elevated temperatures for various strain rates, and the tensile properties of the sheet during warm forming have been evaluated. This study was financially supported by Suzuki Foundation.

5

References

[1] Yoshinaga, H.; Horiuchi, R., Trans. JIM 1963, 4, 134-141. [2] Kaneko, J.; Sugamata, M.; Numa, M.; Nishikawa, Y.; Takada, H., J. Japan Inst. Metals 2000, 64, 141-147. [3] Aida, S.; Tanabe, H.; Sugai, H.; Takano, I.; Ohnuki, H.; Kobayashi, M., J. Japan Inst. Light Metals 2000, 50, 456-461.

469

Stress State Dependent Deformability of Polycrystalline Magnesium T. Törköly, A. Eleöd, Gy. Ziaja Budapest University of Technology and Economics, Budapest

1

Introduction

The twenty-first century has been declared the century of magnesium [1]. It is well known that the mechanical properties of pure magnesium can be substantially improved by alloying or mechanical treatment. As a result of its hexagonal structure mechanical deformability of the pure metal is very limited. The object of this work was to examine experimentally the effect of the stress state on the ductility of pure magnesium and in particular the effect of hydrostatic pressure on deformability.

2

The concept of deformability

Plasticity and deformability are related, but are not identical terms. The plasticity of a given material is interpreted by its resistance to plastic deformation and thus can be characterised by the values of the actual yield stress as a function of deformation, i.e. by the flow curve. The flow stress is dependent on temperature and somewhat on strain rate. Deformability can be defined by the total deformation prior to ductile fracture and is dependent, in contrast to plasticity, on the whole stress state. Maier [2] was the first to plot plastic strain to fracture ε f against the ratio of the average of the three principal stresses and the actual flow stress (kf): (1) ε f = f(k) where

k=

σ m σ1 + σ 2 + σ 3 = kf 3k f

and

kf =

1 2

(σ 1 − σ 2 )2 + (σ 2 − σ 3 )2 + (σ 1 − σ 3 )2

The importance of the Maier’s stress parameter (k) is that it is the ratio of two principal components of the stress state, i.e. it is the quotient of the hydrostatic sphere-tensor and the flow stress, which is connected to the second invariant of the deviator-tensor. Adapting Maier’s representation of deformability, Steiger [3] pointed out that the plastic strain to fracture could not be represented by a single curve but must be depicted by an array of curves. The upper limiting curve is associated with a stress state which is near hydrostatic (σ1 = σ2 > σ3) and the lower one to plane strain conditions (σ2 = (σ1 + σ3)/2), in which the material is less deformable. Bogatov [4] extended the Maier-Steiger representation of deformability by a third invariant of the stress state (i.e. by the Lode-parameter) and represented the relationship of plastic strain to fracture as a function of the stress state indices (k, μσ ) by a fourparametrical function, which describes a surface:

470 ε f = (a + bk + c μσ2 )d

(2)

In Equation (2) the second independent index of the stress state is the Lode-parameter [5]: μσ =

2σ 2 − σ 1 − σ 3 σ −σ = 2 σ1 − σ 3 σ1 − σ 3

(3)

In Equation (2) a, b, c, d are material dependent constants which can be determined experimentally, using four independent investigations. The four-parametrical Bogatov description of the deformability was extended by Darvas and Ziaja [6]. Based on their experimental work, they used an additional function to describe the plastic strain to fracture: ε f = (C1 + C2 με ) exp ⎡⎣(C3 + C4 με )k ⎤⎦

(4)

In Equation (4) C1...C4 are material dependent constants, k the Maier’s stress parameter, and με the Lode-parameter. Darvas and Ziaja [6] suggested the determination of the Lode-parameter by using the strain components obtained from the actual deformation history: με =

2.1

2ε 2 − ε1 − ε 3 ε1 − ε 3

(5)

The influence of hydrostatic pressure on plasticity and deformability

Research on the influence of hydrostatic pressure on different material behaviour goes back over 90 years. Kármán [7] was the first to conclude, using upsetting tests under hydrostatic pressure, that brittleness or plasticity are extrinsic not intrinsic properties of materials, and further that deformability is fundamentally dependent on stress state. The effect of hydrostatic pressure on deformability was studied extensively by Bridgman [8]. In a subsequent monograph Pugh [9] comprehensively reviewed work in the field. He reported that hydrostatic pressure had no effect on the lower yield stress nor on the stressstrain curve. However, a small reduction in the upper yield stress was observed. For more strain hardened specimens the decrease in yield stress was much higher. All researchers agreed that deformability increases under superimposed hydrostatic pressure. Ductile fracture of a free surface always takes place under plane stress condition [6]. On the one hand, we can assume that the presence of the hydrostatic sphere tensor eliminates, or retards the development of the state of plane stress state but on the other hand it can be conducive to the suppression of the formation of voids.

3

Experimental

3.1

Material

The material examined was polycrystalline magnesium, nominally 99.99% pure with a mean grain size of 50 μm. Specimens were machined parallel with the axial direction of the extruded bar, thus retaining in the specimens any preferred orientation in the bar.

471 3.2

Specimens

Table 1 shows the geometry and the stress state coefficients of the test specimens. It is to be noted that all specimens represent different stress states. Table 1. Specimens and test conditions

Investigation

Specimen geometry [mm]

Stress state coefficients k

μσ

-1/3

1

0,41...0,45

1

1/ 3

0

0,72...0,79

1

Upsetting test:

Cylindrical tensile test:

Notched planar tensile test:

Notched cylindrical tensile test:

3.3

Equipment

Upsetting and tensile tests were conducted at room temperature and atmospheric pressure or 400, 600 and 800 MPa hydrostatic pressure. Inside the isostatic pressure chamber, Figure 3, was fitted either a specimen holder for the upsetting tests or a special cartridge containing the specimens for tensile tests. Without filling with oil the pressure chamber could also be used for atmospheric pressure tests. The isostatic pressure chamber was mounted on a electro-hydraulic press. The maximum load on the 40 mm diameter punch was 1000 kN implying a maximum oil pressure of ~800 MPa. During a test the required hydrostatic pressure was created by compressing oil. The last few tenths of a millimetre of punch displacement was used for plastic deformation of the specimens. The amount of deformation was increased by removing oil drop by drop from the pressure chamber. During the tests both force and displacement of the ram were recorded.

472

Figure 3: The isostatic pressure chamber with specimen holder installed for the upsetting test.

4

Results

4.1

Plasticity characterized by the flow curve

Figure 4: Cartridge for tensile tests

The flow curves were plotted from the results of the upsetting tests and are shown in Fig. 5. We can observe that after initial strain hardening a strain softening occurred. It can be in connection with the formation of the slip bands [10].

Figure 5: Flow curves of polycrystalline magnesium

4.2

Deformability values

The strain to fracture was first plotted using the Maier-Steiger limit curves for each test Fig. 6, and subsequently the Bogatov boundary surface was used to define deformability, Fig. 7.

473

Figure 6: Measured deformability of magnesium by different tests

Figure 7: Boundary surface of deformability of magnesium

Fracture of the upset specimens always took place at 45° to the compression axis. Under hydrostatic pressure, instead of fracture, specimens remained deformable (see Fig. 5).

a.

b.

Figure 8: Fracture by upsetting test: a./without hydrostatic pressure, b./ with hydrostatic pressure

Figure 9: Fracture of tensile test

474 The fracture of tensile test specimens was strongly influenced by the actual stress state. With increasing hydrostatic pressure the strain to fracture also increased and specimens became more shifted and twisted with increasing deformation (Fig. 9).

5

Conclusions

Both tensile and compression tests showed that the superimposed hydrostatic pressure component of the stress state markedly increases the deformability of magnesium. In compression tests the decrease in flow stress with increasing strain is associated with the formation of a macroscopic deformation band at 45 degrees to the stress axis in which further deformation is concentrated.

6

Acknowledgements

The authors express their thanks to the Hungarian Research Fund (OTKA T 029963, T 025208) for sponsoring this research, as well as to the CNRS (reference number HJBCU45995) and the Hungarian Academy of Sciences for travel support.

7 [1]

References

Y. Kojima, Platform Science and Technology for Advanced Magnesium Alloys, Materials Science Forum Vols. 350-351 (2000) pp. 3-18. Trans Tech Publications, Switzerland [2] A.F. Maier, Einfluß des Spannungszustandes auf das Formänderungsvermögen der metallischen Werkstoffe. VDI Verlag GmbH, Berlin, 1935. [3] H. Steiger, Über die Abhängigkeit des Formänderungsvermögens metallischer Stoffe vom Spannungszust, Dissertation, TH Aachen, 1965. [4] A.A. Bogatov, O.I. Mizhirickij, V.F. Sisminsev, Ju.A. Aksenov, Fizika metallov I metallovedenie, 45(1978) 5. p.1089-1094. [5] W. Lode, Versuche über den Einfluß der mittleren Hauptspannung aus dem Fließen der Metalle. Zeitschrift für Physik 36(1926) 11/12 p.913-939. [6] Gy. Ziaja, Z. Darvas, Die Abhängigkeit der Umformvermögens Metallischer Werkth stoffe von Formänderungszustand, Metal Forming Theory and Practice, Proc. 5 Int. Conf. on Metal Forming, ICMF’91. Győr, 19-21 June 1991. Ed. M. Tisza, GTE Miskolc, p. 40-47. [7] Th. Kármán, Festigkeitsversuche unter allseitig. Druck, Zeitschrift VDI, 1911, 55, 1749-57. [8] P.W. Bridgman, Studies in Large Plastic Flow and Fracture, McGraw-Hill, London, 1952. [9] H.Li.D. Pugh, Mechanical Behaviour of Materials Under Pressure, Elsevier, 1970. [10] P. Klimanek, A. Pötzsch, Microstructure evolution under compressive plastic deformation of magnesium at different temperatures and strain rates, Materials Science and Engineering A324, Elsevier, (2002) 145-150.

475

A Method to Determine the Damping Behaviour of Coatings on Cast Magnesium Alloy AZ91 J. Göken, D. Letzig, K.U. Kainer GKSS Research Center, Institute for Materials Research, D-21502 Geesthacht, Germany

Abstract Materials damping is a measure of energy dissipated into heat during mechanical vibration under cyclic loading. For magnesium alloys this is mainly due to dislocation motion whereas the damping is reduced by an increasing content of additional elements. Therefore, the commercial cast magnesium alloy AZ91 (Mg-9wt.% Al-1wt.% Zn) which is already established in the automotive industry exhibits relative low damping. Magnesium is a very reactive metal and needs to be protected against corrosion. Selective surface preparation, chemical treatment and coatings are recommended to prevent galvanic attack of magnesium components. Wax coating or painting are possible alternatives for surface treatment of magnesium parts to improve their appearance and corrosion resistance. The influence of wax coating on the strain dependent damping behaviour of AZ91 was studied. The damping was measured as logarithmic decrement of free decaying bending beam vibrations. A strong increase of the damping after coating of the substrate AZ91 was found. Assuming the superimposition principle of damping and a boundary condition with fully adherent phases the logarithmic decrement of the film could be estimated. The results of the damping measurements allow to select a coating material that combines both corrosion protection and damping rise of the substrate.

1

Introduction

Magnesium alloys are the lightest metallic structural materials and attractive in automotive and aerospace industries for saving weight and fuel efficiency, respectively [1, 2]. Although high purity magnesium alloys have improved corrosion resistance galvanic corrosion is still a problem in the presence of electrolytes [3]. Insufficient attention to galvanic corrosion has been one of the major obstacles to the growth of structural applications of magnesium alloys [4]. Therefore, besides the development of suitable chemical compositions special surface treatments are applied to magnesium parts to improve their corrosion resistance. But to provide for a high lifetime of moving parts, control of noise and the stabilization of structures there is also the demand for high damping materials becoming valuable in suppressing critical amplitudes of vibration. The application of high damping construction materials may substitute special energy absorbers or dampers leading to save weight.

476

(a)

(b)

Fig. 1. (a) Time dependence of stress σ and strain ε of a damped material at harmonic excitation leading to a phase lag φ , (b) According hysteresis showing the stored energy W and dissipated energy ΔW .

Damping is the conversion of mechanical energy of a structure into thermal energy. When the material is under cyclic loading anelasticity causes a phase lag φ between stress σ and strain ε (Fig. 1, (a)) and a hysteresis loop is formed (Fig. 1, (b)). The area enclosed by the hysteresis loop represents the energy inside the material during one circle and determines the damping capacity. The specific damping capacity Ψ , the ratio of the dissipated energy ΔW during one complete cycle to the maximum stored energy W from the beginning of the loading to the maximum, is [5] Ψ=

ΔW , W

(1)

with ΔW = 养σ dε

ω t =π 2

and

W=

∫

σ dε .

(2)

ωt =0

It is known that the changes of the microstructure by variation of the chemical composition influence the dislocation damping of AZ91 whereas contrary to such point defect interactions the thermoelastic effect based on the transversal heat flow through a bended material is one of the most important mechanisms leading to a significant increase in damping of this established magnesium alloy at a determined frequency [6]. Because of the superimposition of different damping effects a separation of damping mechanisms is possible since each component contributes to the dissipation of energy and total damping capacity, respectively. Therefore, the damping capacity of the coating material has to be considered and is supposed to be dominant when used as a layer on a low damping substrate to protect it from galvanic corrosion. In this paper the influence of a coating on the damping behaviour of the commercial cast magnesium alloy AZ91 is investigated. The aim is to combine both corrosion protection and increase in damping with suitable coating compounds. It is to be shown that damping measurements facilitate the selection of such coating materials.

477

2

Experimental

For damping measurements the commercial cast magnesium alloy AZ91 (nominal composition Mg-9wt.% Al-1wt.% Zn) was used as received. The sample was machined to a bending beam with a length of 89 mm, a width of 10 mm and a thickness of 3.1 mm. The damping was measured at room temperature in vacuum (60 Pa) as the logarithmic decrement δ of free bending beam vibrations. The bending beam fixed at one end was excited into resonance by means of a permanent magnet screwed at the free end and an exciting coil. The permanent magnet at the beam dipped into a second coil inducing an alternating voltage whose effective value was measured by a multimeter and transmitted to a computer. By measuring simultaneously the effective induced voltage and the amplitudes of various beam vibrations the induced voltage was calibrated in terms of the maximum strain amplitude ε max near the fixed end of the sample (Fig. 2, (a)). The maximum strain amplitude of the bending beam is [7] ε max =

3a ⋅z' . 2l 2

(3)

(b)

(a)

Fig. 2. (a) Bending beam with thickness a , width b and length l moved to z ' at the free end with resulting maximum strain amplitude ε max near the clamped side, (b) Time dependence of ε max of free bending beam vibrations characterizing the materials damping.

For strain amplitude dependent damping the logarithmic decrement is given as δ =

ε 1 ln( max, n ) , k ε max, n + k

(4)

whereas k and n are natural numbers, and k describes the number of cycles between the k-and (n + k)-maximum strain amplitude (Fig. 2, (b)). In case of small damping ( δ Ⰶ 1) [8] follows δ =

1 ΔW . 2 W

Further details about the experimental set-up are described elsewhere [9].

(5)

478 For reasons of simplicity paraffin wax as a coating material was used allowing easy abrasive from and application to the surface of the sample, respectively. In a first step strain dependent measurements of the uncoated magnesium alloy were carried out in the maximum strain amplitude ε max from 10-5 to 10-3. In the following step molten wax was cast on both side of the sample (symmetric trilayer configuration). After cooling down a layer thickness of about 0.7 mm was adjusted by using a heated tool that permitted removal of waste material and closure of possible micro cracks inside the coating compounds by locally liquefaction of the wax. The same damping measurements were repeated with the trilayer bending beam, afterwards with a one-sided coated specimen. The final gaging was performed on the same sample which was completely freed from any surface coating.

3

Results and Discussion

In Fig. 3 the logarithmic decrement δ of the commercial magnesium alloy AZ91 is plotted versus the maximum strain amplitude ε max in logarithmic scale for different coating conditions. 35 30

without wax coating (as received) double-sided wax coated one-sided wax coated wax coating removed

25

δ . 1000

20 15 10 5 0 0.01

0.1

1

εmax 1000 .

Fig. 3. Influence of wax coating on the damping capacity of AZ91.

From former damping measurements it has to be noticed that the damping of the substrate is relatively low in comparison with pure magnesium with about δ = 35 ⋅ 10 -3 at ε max = 1 ⋅ 10 -4 [10]. This can be explained by the amount of foreign atoms in AZ91 working as pinning points and preventing the movement of dislocations that mainly determines the damping behavior of magnesium alloys [11]. A corresponding dislocation string model dealing with the interaction of point defects with dislocations was developed by Granato and Lücke [12].

479 The coating of the substrate generally causes an increase of the damping over the whole strain range of the formed composite material. After double-sided wax coating the damping significantly changed whereas in the strain independent area an increase of damping relative to the bulk material of about eight times occurred. For higher ε max the damping rose leading to a pronounced maximum of the damping observed usually for hysteretic damping in metals. After removal of one surface film this maximum decreased and disappeared after a complete ablation of the coating. The aim is to receive the damping of the film material from the damping of the composite. In Fig. 4 a one-sided coated material is shown. According to the superimposition principle each component of a composite contributes to the dissipation of energy and the materials damping, respectively.

Fig. 4: One-sided coated substrate whose damping capacity is determined by each component.

With Eq. (1) and Eq. (5) the damping of a composite consisting of a substrate and a film material is [13] δC =

1 ΔWS + ΔWF , 2 WS + WF

(6)

whereas the index “ S ” typifies the substrate and “ F ” the film material. Assuming the separation of each damping component for the individual damping follows δS =

1 ΔWS 2 WS

and

δF =

1 ΔWF . 2 WF

(7)

Combining Eq. (6) and Eq. (7) and supposing that there is no damping in the interface between substrate and the film leads to δC =

δ S WS + δ F WF . WS + WF

(8)

In Eq. (8) WS (WS + WF ) and WF (WS + WF ) are weighting factors and represent the relative strain energy fractions of each component of the composite. For small film thickness hF the elastic energy of the film WF can be neglected against WS . With WF Ⰶ WS Eq. (8) can be written as

480 δC = δ S +

WF δF . WS

(9)

For a vibrating bending beam coated one-sided (n = 1) or double-sided (n = 2) with a film material (Fig. 4) with complete adhesion follows [14, 15] δ C = δ S + 3nδ F (

EF hF ), ES hS

(10)

where EF and ES are the elastic modulus of the film and substrate, respectively. Comparing the strain dependent damping curves for n = 2 and n = 1 (Fig. 3) it can be determined that a doubling of the coating with same thickness leads to a mean increase of the composite damping by a factor 2.2 which is in a good agreement with the predicted factor 2 by Eq. (10). From Eq. (10) the film damping can be received as δF =

δ C − δ S ES hS ( ). EF hF 3n

(11)

For many applications the elastic modulus of the used coating material is not available. Knowing both Young’s modulus ES and the specific density ρ , thickness h and resonance frequency f of each composite component the elastic modulus of the film material is [16] EF = (

1 fC2 − fS2 hS ρ F ES . ) + n fS2 hF ρ S 3

Fig. 5. Strain dependent damping of wax calculated by Eq. (11).

(12)

481 In Fig. 5 the strain dependent damping δ F of the film is shown whereas δ F was calculated by Eq. (11). The elastic modulus ES of the substrate was determined as 44 GPa by the measurement of the resonant frequency of the sample [17]. The calculation of EF was performed by Eq. (12) using a density of ρ S = 1.8 g/cm 3 and ρ F = 0.9 g/cm 3 that was taken from the literature [4, 18]. The received value of EF for paraffin wax was 8.3 GPa and within the given range of the elastic modulus of hydrocarbons wax belongs to [19]. Comparing Fig. 3 with Fig. 5 it is evident that in case of a low damping substrate the shape of the composite damping curves is determined by the higher damping component. The damping curves after one-sided or double-sided coating exhibit a similar shape (Fig. 3). For small maximum strain amplitudes the damping is independent from ε max whereas for higher strains a strong dependence on the strain is obvious. Relaxation effects related with the reorientation of atoms by an incorporation of atoms of the coating material into the metal are possible shown for pure iron samples with a synthetic oil film [20]. However, further tests with an oil film on the substrate AZ91 didn’t show any influence of the damping measurements carried out at room temperature and a frequency of about 70 Hz. Therefore, the incorporation of atoms of the coating material into the metal can be disregarded since the damping of AZ91 as received and after complete removal of the layers was not significantly different. Also the dynamic of side chains of the macromolecule wax have to be taken into account [21] for a consideration of energy dissipation of this coating material. But the maximum damping occurs at different positions that would not be expected using the same layer material on both sides of the sample. For metallic foams a significant change of the amplitude dependent damping curves could be observed attributed to both dry sliding of opposite surfaces and energy dissipation in the plastic zones near crack tips [22]. According to these results it is assumed that local incomplete adhesions leading to an additional friction in the interface between the substrate and the film contribute to the measured damping of the composite at higher strains that was neglected for the derivation of the damping of layered materials. The experiments point up the strong damping capacity of wax at an average of about forty times higher than AZ91 whereas in this basic study the mean composite damping leads to a rise in damping of about twelve times relative to the substrate (Fig. 3). The replacement of wax with paintings should lead to similar results but has to be verified.

Summary Materials damping is caused by anelastic effects that can overlap by the superimposition principle. Thus, the damping of a low damping material can be increased by a suitable coating material whose damping capacity must be higher than the substrate. Strain dependent damping measurements of the magnesium alloy AZ91 in terms of the logarithmic decrement were carried out to get information about the damping capacity of the coating material. Assuming a complete adhesion of the substrate and the coating the film damping was calculated by measuring both the substrate and composite damping whereas the thickness and elastic modulus of each component of the composite is necessary to know. The rise in damping after transition from one-sided to symmetrical coating of a bending beam was in a good agreement with the expected doubling of the damping.

482 These measurements allowed to estimate the damping capacity of a coating material whose damping is generally unknown. A systematic coating material selection is optimised including both protection of magnesium against galvanic corrosion and improvement of damping property of the substrate.

References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]

B. L. Mordike, T. Ebert, Mat. Sci. Eng. 2001, A 302, 37-45. E. Aghion, B. Bronfin, Mater. Sci. Forum 2000, 350-351, 19-28. E. Ghali, Mater. Sci. Forum 2000, 350-351, 261-272. M. M. Avedesian, H. Baker, Magnesium and Magnesium Alloys, ASM International, Ohio, 1999. J. Zhang, R. J. Perez, C. R Wong, E. J. Lavernia, Mat. Sci. Eng. Reports 1994, Vol. 13, 8, p. 325. J. Göken, W. Riehemann, Mat. Sci. Eng. 2002, A 324, 134-140. W. Riehemann, Metallische Werkstoffe mit extremer innerer Reibung und deren Messung, Papierflieger-Verlag, Clausthal-Zellerfeld 1994, p. 55. D. W. James, Mat. Sci. Eng. 1969, 4, 1-8. J. Göken, W. Riehemann, Technisches Messen 2001, 68, 535-543. J. Göken, J. Bohlen, D. Letzig, K. U. Kainer in Materials Week 2002-Proceedings (Ed.: Werkstoffwoche-Partnerschaft GbR, Werkstoff-Informationsgesellschaft mbH, Frankfurt, Germany, 2002, p. 627. K. Sugimoto, K. Matsui, T. Okamoto, K. Kishitake, Trans. JIM 1975, 16, 647-655. A. Granato, K. Lücke, Phys. Rev. B 1981, 24, 7007-7017. 3 B. S. Berry in M D: Mechanics and Mechanisms of Material Damping (Ed.: V. K. Kinra, A. Wolfenden), ASTM, Baltimore, USA, 1992, p. 28. M. Weller in Mechanical Spectroscopy Q-1 2001 (Eds.: R. Schaller, G. Fantozzi, G. Gremaud), Trans Tech Publications Ltd, Zürich, Switzerland, 2001, Chapter 8. S. I. Hayek, Encyclopedia of Applied Physics (VCH Publishers Inc.) 1994 Vol. 9, p. 561. B. S. Berry in Diffusion Phenomena in Thin Films and Microelectronic Materials (Ed.: D. Gupta, P. S. Ho), Noyes, Park Ridge, NJ, 1988, p. 43. A. S. Nowick, B. S. Berry, Anelastic Relaxation in Crystalline Solids, Academic Press, New York, 1972, p. 630. G. W. Frasier, J. Keiser, J. Range Manage. 1993, 46, 91-93. M. F. Ashby, Materials Selection in Mechanical Design, 2nd ed., ButterworthHeinemann, Oxford 1999, p. 48. D. Samatowicz, J. Alloys Comp. 2000, 310, 457-461. -1 J. Y. Cavaillé, L. David, J. Perez in Mechanical Spectroscopy Q 2001 (Eds.: R. Schaller, G. Fantozzi, G. Gremaud), Trans Tech Publications Ltd, Zürich, Switzerland, 2001, Chapter 7. I. S. Golovin, H.-R. Sinning, J. Göken, W. Riehemann, Solid State Phen. 2003, 89, 267-272.

483

Microstructure and Mechanical Properties of Heat Treated Mg-Zn-Al Based Alloys Valérie Lemaire, Dominique Dubé, and Réal Tremblay Université Laval, Québec, Canada

1

Abstract

The microstructure of cast magnesium-zinc-aluminum based alloys can be modified by heat treatment and their mechanical properties can be improved at room temperature as well as elevated temperature. In this study, specific heat treatments have been performed on ZA104 and ZA10403 (0,3%Ca) alloys in order to precipitate intermetallic phases in the matrix and determine their effects on mechanical properties. The hardness, tensile strength, ductility, and creep resistance of these alloys has been determined as a function of temperature and duration of heat treatment. The evolution of microstructure with heat treatment is reported and discussed in order to explain the changes in mechanical behaviour.

2

Introduction

In the last ten years, many new families of alloys have been developed for applications where creep resistance is necessary. The performances of the new alloys have been evaluated and significant progress has been achieved, most alloys being developed with the objective to reach good castability, sufficient strength and ductility at room temperature, better creep strength, and the lowest cost. Among the creep resistant alloys, studies have demonstrated that ZA (Mg-Zn-Al) and ZAC (Mg-Zn-Al-Ca) alloys containing 8-12%Zn and 4-6%Al display a significantly improved resistance to creep [1-8], the addition of calcium enhancing their resistance. Recent works also showed that the microstructure of Mg-Zn-Al based alloys can be modified substantially by heat treatments [4,7,9,10] and mechanical properties improved. More recently, Mendis et al. [8] and Vogel et al. [9] have identified the crystal structure of precipitates modifying the modification in mechanical properties of Mg-8Zn-4Al and Mg8Zn-5Al alloys, showing evidence of quasi-crystalline phases. Two main aspects must be investigated in order to obtain an effective heat treatment that best improves the mechanical properties of Mg-Zn-Al based alloys. Firstly, the composition of an alloy that offers the best response to a specific heat treatment must be determined. Secondly, the parameters oh the most effective heat treatment must be established. It involves the consideration of parameters such as the maximum temperature and time cycle, based not only on knowledge of phases formed during aging but also on their influence on mechanical behaviour.

484 In the present work, two alloys have been selected among a group of Mg-Zn-Al based alloys recently developed (Mg-10Zn-4Al and Mg-10Zn-4Al-0,3Ca) in order to obtain sufficient information about the influence of heat treatment on various mechanical properties. Selected heat treatments were carried out with the objective to vary the microstructure of these alloys and improve their performance.

3

Experimental Procedure

ZA104 alloy has been prepared by melting pure magnesium, zinc, and aluminum in a stainless steel crucible. In order to determine the influence of calcium on aging and mechanical properties, calcium from Mg-Ca master alloy was added to the melt and specimens of ZAC10403 alloy have been produced. Molten alloys have been poured in stainless steel moulds coated with appropriate refractory. Flat specimens (5 mm thick) have been cast for hardness measurements. Round specimens (12,7 mm diameter) were cast according to the ASTM standard E8 for tensile and creep properties evaluation. Aging response was obtained by heat treatment in an electrically heated furnace, the maximum temperature selected ranging between 150 °C and 320 °C. The temperature of heat treatment was selected following indications from cooling curves [6,7]. The duration of treatment at maximum temperature varied between 1 h and 100 h and rapid cooling was applied at the end of thermal cycles. -1 Mechanical properties evaluated include tensile tests at 0,0016 s , creep tests at 34,5 MPa, 175 °C ± 1 °C and 200 h , and hardness tests with Rockwell E (HRE) scale. Microstructure was examined using scanning electron microscopy (SEM) and optical microscopy (OM) and analysis of phases was performed by electron probe microanalysis (EPMA).

4

Results and Discussion

4.1

Microstructure in as-cast condition

Figure 1 shows the microstructure of ZA104 and ZAC10403 alloys in as-cast condition, revealing the magnesium-rich dendrites surrounded by microconstituent containing intermetallic phases of structure similar to that of Mg32(Al,Zn)49, τ1' and τ2' [7,8,11]. The new τ2' phase which contains more calcium also provides much higher thermal stability [7,8]. In recent structural analysis, the crystalline structure of intermetallic Mg32(Al,Zn)49 has been presented as quasi-crystalline [8,9]. No trace of MgZn (ε) or Mg5Zn2Al2 (φ) phase was found in as-cast ZA104 based alloys [7].

485

(a)

(b)

Figure 1. Microstructures of ZA104 (a) and ZAC10403 (b) alloys in as-cast condition.

4.2

Microstructure after heat treatment

Figure 2 shows that after 100 h heat treatment at 320°C, the microstructure of Mg-10Zn4Al based alloys is not homogeneous. A great proportion of the intermetallic phases, isomorphous with Mg32(Al,Zn)49 formed during solidification, remains after heat treatment. Clusters of large precipitates (particles) are also found and for the same treatment, no precipitate is visible near the interface of intermetallic phase with the matrix (Figure 3).

Figure 2. Typical microstructure of Mg-10Zn-4Al alloy after heating 100 h at 320 °C showing a smoothening of intermetallic phases and large particles (arrow).

Figure 3. Interface of intermetallic phase with α-Mg matrix after a treatment of 100 h at 320 °C. No precipitate is visible.

For specimens heat treated at 200 °C for 137 h, a precipitation free zone (PFZ) and a band of precipitates are visible in the vicinity of τ1' and τ2' (Figure 4). No crystallographic study was performed on these precipitates that have been recently identified as quasicrystalline phase [9,10]. The presence of a supersaturated zone surrounding intermetallics enhance precipitation [3,9]. The use of SEM allowed observing the presence of fine precipitates only in specimens aged at temperature above 200 °C, but they are suspected in specimens aged at 150 °C and 170 °C as confirmed later by changes in mechanical properties. However, a solution treatment at 320 °C for 100 h followed by aging at 170 °C for 24 h has permitted to show the precipitation in ZA104 alloy (Figure 5). The absence of fine precipitates after treatment at 320 °C indicates above solvus conditions and the microstructure confirms that the specimens were treated below the temperature of the last solidification event (Figure 3).

486

Figure 4. Fine precipitates typically found in ZA104 and ZAC10403 alloys after aging at 200 °C for 137 h.

4.3

Figure 5. Microstructure of ZA104 alloy after solution treatment of 100 h at 320 °C followed by aging at 170 °C for 24 h.

Hardness curves

Figure 6 shows the aging response of ZA104 and ZAC10403 alloys heat treated at various temperatures. For treatments carried out at 150 °C and 170 °C, the hardness increases rapidly and reaches a maximum value after 24 to 48 h, being stable at about 86-88 HRE. The increase in hardness is moderate for treatments performed at 200 °C. Heat treating at 320 °C slightly decreases the hardness. The homogenization of aluminum and zinc within primary α-Mg dendrites reported previously reduce supersaturation thereby explaining the observed decreasing in strength [3].

(a)

(b)

Figure 6. Aging curves at different temperatures for ZA104 (a) and ZAC10403 (b) alloys.

Based on these results, combined heat treatments were performed, including prior solution treatment at 320 °C and aging at 170 °C which had given the highest hardness. Figure 7 shows the hardness as a function of time for different heat treated specimens. ZA104 specimens (Figure 7a) heat treated at 170 °C showed rapid aging and stabilization of hardness at 87 HRE whereas specimens heat treated at 320 °C present a decrease in hardness due to homogenization of solutes Al and Zn within dendrites. The combination of the partial solution treatment at 320 °C followed by an aging treatment at 170 °C brought a slower aging and a lower hardness was reached (84 HRE) after 80 h. Similar results were observed for ZAC10403 heat treated specimens (Figure 7b), meaning that such content in calcium has no significant effect on aging.

487

(a)

(b)

Figure 7. Aging curves showing the influence of combined treatments for ZA104 (a) and ZAC10403 (b) alloys.

4.4

Tensile and creep properties

Table 1 shows the mechanical properties, strength, ductility and creep behaviour as a function of parameters used in some heat treatments for ZA104 alloy. Properties of AZ91D alloy are also indicated as a reference. Table 1. Tensile and creep properties for ZA104 alloy after different heat treatments. Properties for AZ91D alloy are also given as a reference.

Alloy

Heat treatment

Tensile Y.S. (MPa)

Elongation (%)

ZA104

as-cast 24 h / 170°C 100 h / 320°C 100 h / 320°C + 24 h / 170°C as-cast

183 138 82 127 160

AZ91 (1)

2,6 1,6 2,9 1,0

Total creep strain(1) (%) 0,433 0,223 0,808 0,880

Min. creep rate(1) (%/h) 4,5 x 10-4 2,8 x 10-4 10,0 x 10-4 -4 22,9 x 10

3,0

1,33

24,0 x 10-4

At σ = 34,5 MPa, T = 175 °C and t = 200 h.

The best combination of tensile strength and ductility is obtained in as-cast condition. However a significant improvement in creep is reached after aging at 170 °C. The partially homogenized bars at 320 °C did not produce improved strength at room temperature or improved creep probably due to the loss of supersaturated zone near intermetallic phase that have crystallized during solidification. It becomes evident that precipitation takes place during creep at 175 °C. A stabilization treatment such as aging contributes to decrease the plasticity of the alloy but improves creep resistance. However the best combination of temperature and time remains to be accomplished to get the optimum results.

488

5

Conclusions

The present investigation on ZA104 and ZAC10403 alloys leads to the following conclusions: 1. The microstructure of the two ZA alloys can be modified by appropriate heat treatments. 2. The tensile strength and hardness are increased but ductility is reduced. 3. Creep behaviour is best improved by aging at 170°C. 4. Calcium addition below 0,3% has not strongly affected the maximum hardness obtained by heat treatment. 5. There is no direct correlation between hardness and creep behaviour.

6

Acknowledgements

The authors would like to thank the NSERC (Canada) and FQRNT (Québec) for financial support, Norsk-Hydro (Québec) for magnesium supply, M. Larouche, M. Choquette and G. Bureau for technical assistance.

7

References

[1] Z. Zhang, A. Couture and A. Luo, Magnesium Alloys and their Applications (Ed. B.L. Mordike and K.U. Kainer), Wiley-VCH, Weinheim, Germany, 1998, 289 - 294. [2] Z. Zhang, A. Couture and A. Luo, Scripta Materialia, 1998, 39, 45 - 53. [3] Z. Zhang, Development of Magnesium-Based Alloys for Elevated Temperature Applications, Ph.D. Thesis, Univ. Laval, Québec, Canada, 2000, 205 p. [4] I.A. Anyanwu, S. Kamado, T. Honda, Y. Kojima, S. Takeda and T. Ishida, Mater. Sc. Forum, 2000, 350-351, 73 - 78. [5] Z. Zhang, R. Tremblay, D. Dubé and A. Couture, Can. Metal. Quart., 2000, 39, 503 512. [6] Z. Zhang, R. Tremblay and D. Dubé, Magnesium Technology 2001 (Ed. J. Hryn), TMS, 2001, 147 - 151. [7] Z. Zhang, R. Tremblay and D. Dubé, Mater. Sc. and Tech., 2002, 18, 433 - 437. [8] C.L. Mendis, L. Bourgeois, B. Muddle, and J-F Nie, Magnesium Technology 2003 (Ed. H.I. Kaplan), TMS, 2003, 183 - 188. [9] M. Vogel, O. Kraft, G. Dehm and E. Artz, Scripta Materialia, 2001, 45, 517 - 514. [10] C.L. Mendis, Precipitation Hardening of Magnesium-Zinc-Aluminum Casting Alloys, Master Thesis, Monash Univ., Clayton, Australia, 2000, 139 p. [11] Z. Zhang, R. Tremblay, D. Dubé and A. Couture, Proc. Inter. Symp. on Light Metals th 2000 (Ed. J. Kazadi and J. Masounave), 39 Conf. of Metallurgists of CIM, Ottawa, Canada, 2000, 273 - 284.

489

Analysis of Deformation and Fracture Mechanisms in Mg-Az31 Sheet Material Gaofeng Quan, Jürgen Heerens Institute for Materials Research, GKSS Research Centre, 21502 Geesthacht, Germany

Abstract In the present work, deformation and fracture in rolled Mg AZ31 sheet material are investigated using tensile specimens with different orientations, tested at various strain rates. In RD specimens (loading in rolling direction) the generation of dislocation slip bands is predominately observed. This correlates with a low strength of the material and small plastic anisotropy value, r. In TD specimens (loading in transverse direction) the material deforms due to intense twin development, which is combined with a high strength and a large r-value. The r-values increase with increasing strains but seem to be independent on the strain rate. In RD specimens micro-cracks nucleate predominantly at dislocation slipbands. In TD specimens the micro cracks initiate mainly at intersections and interfaces of twin bands. In both orientations a slant fracture surface is observed on macroscopic scale.

Introduction Correlations between microstructure and macroscopic mechanical properties of metallic materials are of particular interest in several research areas. Typical areas are mechanical property enhancement or development of numerical models for simulation of material properties. For Mg-alloys, as a relatively new prospective material used for lightweight structures, the knowledge about the mechanisms of deformation and fracture on microscopic and macroscopic scale is still very limited. Mg-alloys have an hcp crystal lattice, and therefore their mechanical properties are different compared with these of Al-alloys. In general, cast and wrought Mg alloys have strong textures [1, 2, 3] and hence, anisotropic mechanical properties [4, 5]. Mg-AZ31 is the most popular wrought Mg-alloy, usually offered as sheet material with small thickness. This type of alloy is still under investigation in order to improve its mechanical properties as required for deep drawing and other conventional forming routines. Its anisotropic mechanical properties and its low plastic failure stain are main obstacles for the practical application of this alloy.

490 Several published investigations deal with microstructure evolution and creep property at elevated temperature but there are scarce results on the deformation, damage and fracture properties of AZ31 sheet material [6]. In the present paper the effects of orientation and strain rate on the macroscopic and microscopic deformation and fracture behaviour of Mg AZ31 sheet material are investigated. The aim is to provide basic knowledge and test data as required for the development and validation of micro-mechanical models for simulating the mechanical properties [7, 8]. The measured results of the texture in this material appears also on this conference [9], which shows strong fibre texture in (0002) pole figure, i.e., the (0001) lattice planes prefer to be parallel to rolling plane, that is one of the most important anisotropic factor of the microstructure. Based on these results, the aim of this paper is to investigate the mechanisms of plastic deformation, crack nucleation and fracture by tensile tests.

Experimental procedures Material and specimens The material used in the present paper is a hot rolled AZ31 Mg alloy, with a thickness of 1.6 mm. From the sheet material tensile specimens were machined. The specimen width was 20 mm, the gauge length 50 mm and the thickness 1.6 mm. Some of the specimens were polished and etched to observe the surface morphology before and after tensile tests. Specimens were prepared in three orientations: rolling direction (RD), 45º to RD and transverse direction (TD). In this experiment, L indicates the loading direction, T the transverse or lateral direction to the loading direction, and S the sheet thickness direction. Tensile Tests The tests were performed using a servo-hydraulic 25 kN testing machine. The tensile tests were conducted at 3 strain rates: 1.1 × 10-4, 1.1 × 10-3, 1.1 × 10-2 s-1 achieved by crosshead speeds of 0.4 mm/min, 4.0 mm/min, and 40 mm/min. The specimen elongation and width reduction were measured using several clip gauges and a video system, and the reduction in thickness was measured after fracture.

Results and discussion The specimen orientation and the strain rate have a strong affect on both proof stress and ultimate strength. The more the specimen deviates from the rolling direction, and the more the strain rate increases, the higher the strength of the material is, see Fig.1. Deformation data measured at the specimens were used to derive the plastic anisotropy value, r, see Fig.2. It is seen that the r-value depends on the strain level but not on strain rate. This tendency is strong in TD orientation but small in RD orientation.

491

Fig1. Effects of strain rate and orientation on yield strength

Fig.2. Effects of strain, strain rate and specimen orientation on r-value

Effects of specimen orientation and strain rate on the ductility are outlined in Fig. 3a, b. The diagrams show the strains in S direction, in T direction and in L direction at final fracture. The T and S strains have been measured close to the fracture surface within the necking regime. The strains have the following main characteristics: • With increasing strain rates the L-strain and T-strain decrease, whereas the S-strain is less affected, Fig.3a. In TD specimens the change is more obvious. • With increasing deviation of the specimen orientation from the rolling direction, the L and T-strain increase, whereas the S-strain decreases, Fig.3b. At higher strain rate the tendency is depressed.

Fig. 3a Strain rate vs fracture strains in 3 specimen orientations.

Fig. 3b Specimen orientation vs fracture strain at 3 different strain rates.

492

Damage and fracture observation Micro crack nucleation In the as-rolled condition, the material shows twins or twinning bands located within the grains, see Fig. 4. The twins and the twin bands are randomly distributed in the material. Twins and twinned bands always end at grain boundaries. a)

b)

Fig. 4. Original twin morphology on LT plane: a) optical micrograph and b) SEM image

It is known that at room temperature plastic deformation in this material is predominantly due to dislocation slip in the ((0001)< 1120 >) basal-plane system and due to primary twinning, which occurs at the ( 10 12 ) and ( 3034 ) planes. Grain rotation, particle (inclusions) facture are less decisive in the micro-crack nucleation processes. Fig. 5 shows surface morphologies of polished RD and TD specimens close to the fracture surface. Comparing the two micrographs it is seen that in TD specimens, the twinning mechanism is dominant whereas in RD specimens dislocation slipping is the most active process. a)

RD and Loading direction

Fig. 5a. SEM surface morphology of polished RD-tensile specimen after fracture. (Image is parallel to the rolling plane)

493 b)

Loading direction RD

Fig. 5b. SEM surface morphology of polished TD tensile specimen after fracture. (Image is parallel to the rolling plane)

On both surfaces a lot of micro cracks are seen. In TD specimens the micro cracks initiate at the borders or intersections of twinned bands, see Fig. 5b. In RD-specimens the micro-cracks generate perpendicularly to the slip bands. Fractography For all specimen orientations the macroscopic fracture surface is inclined to the loading direction and thickness direction by about 45°, but laterally the fracture edge is normal to loading direction. Fig. 6a shows a typical example of SEM micrograph obtained from a RD-specimen. The fracture surface shows some steps. On the steps one can find flat areas, which have a scale of 20~50 μm, and similar to that of the grain size. These areas show traces of plastic deformation as seen in Fig.6b. Broken particles or cavity growth around particles were not found on more than 10 fracture surfaces investigated. a)

b) 45°

Fig. 6. Fracture surface morphology of tensile specimens. a) RD, general morphology and steps parallel to the free surface, b) RD, local image in the square of a), details in a flat area. (Micrographs are parallel to the macroscopic crack plane.)

494

Summary The AZ31 sheet material shows strong anisotropy in strength, ductility, deformation and fracture mechanisms, which can be summarised as follows: 1. In TD specimens the strength and ductility are higher than in RD specimens. 2. The anisotropy factor of plastic deformation, r, is large in TD specimens and small in RD specimens. No significant strain rate dependence of r was found. 3. In RD specimens plastic deformation occurs predominantly via dislocation slip-bands. In TD specimens intense twinning is observed. 4. In RD specimens, micro-cracks nucleate at the dislocation slip-bands. The micro cracks propagate perpendicular to the slip bands and 45° to the loading direction. 5. In TD specimens, micro-crack nucleation occurs at the intersections and interfaces of the twinned bands. 6. In TD and RD specimens the macroscopic fracture edge is oriented 90° to the loading direction and inclined by 45° in thickness direction (slant fracture).

Acknowledgement Many thanks are due to Kay Erdmann for his support in performing the tests. Thanks are also due to Dr. D. Letzig and Dr. J. Bohlen for their helpful discussion and for providing the test material.

References [1] Hydro Magnesium, Data Sheet, Die Cast Magnesium Alloys, 1999 [2] K. U. Kainer, F. von Buch, Mat.-wiss. u. Werkstofftech, 1999, 30, 159-167 [3] USGS (U. S. Geological Survey), Magnesium: Statistics and Information, 2001 and 2002 [4] E. Aghion, B. Bronfin, D. Eliezer, J. Mater. Proc. Techn., 2001, 117, 381-385 [5] W. J. Kim, C. W. An, Y. S. Kim, S. I. Hong, Scripta Materialia, 2002, 47, 39-44 [6] M. M. Myshlyaev, H. J. McQueen, A. Mwembela, E. Konopleva, Mater. Sci. Engng, 2002, A337, 121-133 [7] O. Chabanet, D. Steglich, J. Besson, V. Heitmann, D. Hellmann, W. Brocks, Computational Material Science, to be published, 2003 [8] J. Heerens, D. Steglich, W. Brocks, ABAQUS User Conference, Freiburg, Germany, 2001 th [9] F. Kaiser, J. Bohlen, D. letzig, K. U. Kainer et al, 6 International Conference on Magnesium, alloys and their applications, Wolfsburg, Germany, Sept. 2003

495

Unstable Plastic Deformation in Mg Alloys-post Relaxation Effect Zuzanka Trojanová1, Pavel Lukáč1, Luboš Čížek2 1 2

Faculty Mathematics and Physics, Charles University, Ke Karlovu 5, CZ-121 16 Praha 2 Technical University Ostrava, Tř. 17. listopadu 2172, CZ-708 00 Ostrava

Abstract Post relaxation effect, i.e. an increase of the flow stress at the beginning of plastic deformation after stress relaxation in comparison to the flow stress at the beginning of stress relaxation, is observed in the case of dynamic strain ageing. The post relaxation effect is sensitive to the strain and temperature. Possible changes in the internal stresses and the mobile dislocation density are considered to be responsible for the post relaxation effect.

Introduction The plastic deformation of alloys exhibit many phenomena associated with mobile solutes. Following features give an evidence of dynamic strain ageing (DSA): (1) a local maximum in the temperature dependence of the yield stress; (2) a local maximum in the temperature dependence of the stress sensitivity parameter n; (3) a stress increase after stress relaxation in the DSA temperature interval. The observed anomalies are caused by the change in the concentration solute atoms on dislocations with the strain rate and temperature. The stress relaxation (SR) technique has been demonstrated to be quite useful experimental method. The specimen is deformed to a certain stress σ0 and then the machine is stopped and the stress decreases with time t. The specimen can be again reloaded to a higher stress (load) and the test may be repeated. The time derivative σ& = dσ/dt is the stress relaxation rate and σ = σ(t) is the flow stress at time t during the SR. Stress relaxation is very often analysed under the assumption that the stress relaxation rate is proportional to the strain rate ε& according to [1] ε& = −σ& / M,

(1)

where M is the combined modulus of the specimen – machine set. In order to estimate parameters of the thermally activated processes from the SR the following theoretical relationship between stress drops rate σ& and stress σ can be used [1] ln ( −σ& ) = C + n ln σ

(2)

496 where C is a constant and n is the stress sensitivity parameter defined as ⎛ d ln ε& ⎞ n=⎜ ⎟ . ⎝ d ln σ ⎠ T

(3)

The SR are usually analysed assuming constancy of the mobile dislocation density ρm and internal stress σi. The stability of structure during the SR is a main condition for the correct determination of the activation volume and consequently for clearing up the microscopic mechanism controlling the dislocation mobility in a material. An unstable structure may influence the course of the SR. One of the reason for the structure instability (the dynamic strain ageing phenomena) in an alloy is the solute atoms diffusion towards moving dislocations during deformation: the Portevin-Le Chatelier effect, the positive dependence of the flow stress on temperature and its negative dependence on strain rate, the sharp decrease in the creep rate. The other macroscopic indication of the dynamic strain ageing is a post relaxation effect - the pronounced yield point appearing at sample reloading after the SR. The purpose of this work is to analyse SR tests performed in three magnesium alloy in the dynamic strain ageing temperature interval.

Experimental Procedure Two magnesium alloys were used in this study. Magnesium alloy AZ91 (nominal composition 9 wt% Al, 1 wt% Zn, 0.3wt% Mn, rem. Mg) was prepared by the squeeze cast technology. Thermal treatment T6 was used according to Polmear [2] (homogenisation for 18 h at 413 °C, then precipitation for 8 h at 168 °C). Magnesium alloy WE54 (nominal composition 5 wt.% Y, 2 wt% Nd, 2 wt% mixture of rare earth – Tb, Er, Dy, Gd, rem Mg) was prepared by powder metallurgical technology and extruded at elevated temperature. Figure 1: Post relaxation effect-schematic drawn.

Samples were deformed in an INSTRON 1195 machine in the temperature interval of 22-300 °C. Sequential stress relaxation tests were performed at increasing vs tress along a stress-strain curve. Duration of the SR was 300 s. Deformations at room temperature were -5 -4 -1 performed at various initial strain rates in the interval from 2.2 x 10 to 6.7 x 10 s and duration of the SR from 120 s to1500 s. SR were analysed according the equation (2). The slope of the plot ln(- σ& ) vs ln(σ/σ0) (σ0 is the starting stress of the SR) is the stress sensitivity parameter n defined by (3). As the characteristics of the post relaxation effect, the stress increment Δσ after reloading of the sample was used.

497

Experimental Results and Discussion

ln(-σ) [MPa/s]

Examples of the SR curves in a ln(- σ& ) vs ln(σ/σ0) representation obtained for WE54 alloy are introduced in Fig. 2 (for room temperature, duration of 600 s and various starting stresses of the stress relaxation test. Tem6 -5 -1 2.2x10 s perature dependence of the stress sensitivity 278 317 5 parameter n estimated for the first SR (in 349 the vicinity of the yield stress) and for the 420 4 last SR (in the vicinity of the maximum stress) are shown in Fig.3. The introduced 3 dependencies show a local maximum at 50 °C which indicates a presence of the 2 dynamic strain ageing process. Stress sen1 sitivity parameter obtained for WE54 alloy 0.96 0.97 0.98 0.99 1.00 at various starting stain rates depending on the stress is introduced in Figure 4. σ/σ0 The dynamic strain ageing is the pheFigure 2: SR curves obtained for various starting nomenon of interactions between moving stresses at room temperature for WE54 alloy dislocations and diffusing solute atoms. The impurity diffusion towards both immobile and mobile dislocations has been regarded in many works [3-5]. At size interaction impurities segregate at small times t, according to the law ⎛ Dt ⎞ cd = c0 ⎜ α 2 ⎟ ⎝ b ⎠

2/3

(5)

where cd is the solute atoms concentration, D is the diffusion coefficient for the solute atoms and α is a constant (depending on the interaction between a solute atom and the dislocation) and b has a magnitude of the lattice constant. The model developed by authors of the paper [6] considers that the dislocations are not waiting passively for the solutes to diffuse to them, but rather they capture the solutes during their bowing out process, under the applied stress, while being held at dislocation forest. The increase in the concentration of the dislocation core atmosphere is a result of the diffusion of the existing core atmosphere along with the dislocation, and the new solutes which are encountered and captured by the dislocation. Due to the increase in the core atmosphere density, both the total thermal activation energy and the mechanical threshold stress are increased. In an alloy the flow stress may be consider as a sum of two additive contributions: σ = σf + σd ,

(6)

with σf relating to a friction imposed by the solutes on the mobile dislocations, and σd relating to the dislocation-dislocation interaction.

498 400

240 200

σ02

σmax

200

n

n

300

-5 -1

2.2x10 s 5.6x10-5s-1 -4 -1 5.6x10 s

160 120

100

80

0 0

100

200

300

100

200

300

400

500

σ0 [MPa]

T [°C]

Figure 3: Temperature dependence of the stress sensitivity parameter obtained for AZ91 alloy.

Figure 4: Influence of the strain rate on the stress sensitivity parameter obtained for WE54 alloy.

The strain rate sensitivity of the flow stress has also two additive contributions: dσ f dσ = d ln ε& d ln ε&

+

d ln σ d d ln ε&

(7)

The strain rate sensitivity of the flow stress is done as

dσ d ln σ = σ = σ f m f + σ d md d ln ε& d ln ε&

(8)

Stress sensitivity parameter measured in the SR test is n = d ln ( −σ& ) /dlnσ = dln ε& /dlnσ n =

d ln ( −σ& ) d ln σ

=

σ f mf

σ . + σ d md

(9)

The strain rate sensitivity component mf has in the dynamic strain ageing regime an anomalous negative value [7] and then parameter n may have a local maximum in the temperature dependence of this parameter. Such maximum can be seen in Figure 3. As it is to see from Figure 4 parameter n depends in the DSA region of the strain rate. We can assume that during the stress relaxation test solute atoms diffuse to dislocations. Dislocations are pinned by these solute atoms. Therefore an increase in the stress is needed to move dislocations after reloading, i.e. the flow stress at the start of new deformation after stress relaxation is higher. The stress increment after reloading of the sample Δσ was measured for AZ91 alloy depending on the time for at starting stresses of the SR (Figure 5). The dependence of Δσ on the relaxation duration is well given by a power law with the slope s = 0.2, which is lower value in comparison with the prediction of the theory.

499 8

50°C 100°C 150°C reg.

10

σ0=157 MPa

Δσ [MPa]

Δσ/MPa

6 4 2

σ0=175 MPa reg. line

0

1 2

5

10

0

25

1

2

3

ε [%]

time/min

Figure 5: The time dependence of Δσ estimated for AZ91 alloy at room temperature and two starting stresses of the SR test.

Figure 6: Strain dependence of the stress increment obtained for AZ91 alloy and three temperatures.

The strain dependence of the stress increase Δσ was measured for AZ91 alloy and three temperatures (Figure 6) and for WE54 alloy at room temperature and three strain rates (Figure 7). The stress increment after SR is given by the relation [8]: Δσ(t,ε,T) = Δσm(ε,T) {1 – exp[-(t/tc)p]} ,

(10)

where Δσm(ε,T), the stress increment for t → ∞, grows with increasing size misfit of the solute and matrix atoms, increasing solute atoms concentration and lowering temperature. 16

5 Δσ [MPa]

12 Δσ [MPa]

6

2.2x10-5 5.6x10-5 5.6x10-4

8 4

4 ε = 0.5% ε = 1.5%

3

0 0

2

4

6

8

10

12

14

16

ε [%]

Figure 7: Strain dependence of Δσ for WE54 alloy -1 and three strain rates in s .

2 0

40

80

120

160

T [°C]

Figure 8: Temperature dependence of Δ obtained for AZ91 alloy.

tc is a characteristic time which depends on the strain as tc ~ ε-k . Power coefficient k was found to be for In based alloys k=1-2 [9]. Strain and temperature influence significantly the Δσ values as it can be seen from Figures 7 and 8. It is reasonable to expect that dislocation from more than one slip system form dislocation pile-ups. The strain ageing is due to migration of dissolved solute atoms towards the pile-ups. The total dislocation density is increasing with increasing strain while the number of the free solute atoms is constant. Therefore the strength of dislocation locking is smaller for higher strain degree. On the other hand Dlouhý, et al. [10] have shown that changes in the internal stress and the density of moving dislocations with the time of SR can cause the stress increment after the reloading. The stress increment is a complex function of the stress relaxation time.

500

Conclusions (1) Stress relaxation curves of Mg alloys AZ91 and WE54 were studied in the stress rep& ) vs ln σ plots. resentation i.e. ln(- σ (2) Stress sensitivity parameter n depends on the temperature and the strain rate. (3) A stress increase after stress relaxation was observed. (4) Observed post relaxation effect as well as anomalous temperature dependence of n are caused by the change in the concentration solute atoms on dislocations with the strain rate and temperature.

Acnowledgements The authors thank the Grant Agency of the Academy of Sciences of the Czech Republic for financial support under grant A2112901.

References [1] V.I. Dotsenko: phys. stat. sol. (b) 1979, 93, 11. [2] I.J. Polmear: Magnesium Alloys and Their Applications (Eds. B.L. Mordike and F. Hehmann) DGM, Oberursel, 1992, 201. [3] A.H. Cottrell, M.A. Jawson: Proc. Roy. Soc. 1949, A199, 104. [4] H. Yoshinaga, S. Morozuni: Phil. Mag. 1071, 23, 1367. [5] G.A. Malygin: phys. stat. sol. (a) 1982, 72, 493. [6] J. Cheng, S. Nemat-Nasser: Acta mater. 2000, 48, 3131. [7] S.I. Hong: Mater. Sci. Engn. 1085, 76, 77. [8] S.V. Lubenets, L.S. Fomenko: Fizika metallov metall. 1986, 62, 377. [9] S.V. Lubenets, V.I. Startsev, L.S. Fomenko: Czech. J. Phys. 1986, B36. 493. [10] A. Dlouhý, P. Lukáč, Z. Trojanová: Kovové mater.1984, 26, 688.

501

On the Reliability of Yield Strength Data Evaluated from as-cast Tensile Test Bars of Mg-alloys M. Wessén, H. Cao Jönköping University, Jönköping, Sweden

1

Introduction

Within the automotive sector, energetic efforts are today being made to decrease vehicle weight in order to minimize the emission of combustion gases. In the light of this work it is natural that the use of light metals, and especially magnesium, has increased considerably during the last 5-10 years. Some applications in cars where cast magnesium components are used today include seat frames, instrument panels, transfer cases, wheels, steering wheels and various kinds of housings. During the optimization process of new components, stress/strain analysis tools are usually used to calculate the stress state when the component is exposed to its working loads. In order to avoid any permanent plastic deformation it is of course necessary that the stress doesn’t exceed the local yield strength in any part of the component. It should also be emphasized that cast Mg-alloy components can have a very large variation in yield strength in different regions as a result of the varying cooling conditions and solidification microstructures [1-2]. Some materials have a distinctly defined yield point, e.g. most steels, whereas other materials start to plastify gradually at very low stresses. In the latter case the yield strength is most commonly defined as the stress when a permanent deformation of 0.2% remains after unloading; this is also often called the 0.2% proof stress. The plastic strain that occurs prior to the proof stress is often called “pre-yield micro strain”. Magnesium alloys usually exhibit this behaviour. As a consequence, it is difficult to define the Young’s modulus (E) in Mg-alloys, since E varies with stress at even moderate stresses. The elastic-plastic transition and the Young’s modulus in magnesium alloys AZ91 and AM60 (heat-treated samples) have been studied thoroughly by Cáceres et al [3]. When studying published stress-strain curves for cast Mg-alloys it seems that samples cast in sand moulds as well as heat treated samples do not show a well defined yield point (as discussed above), and that the curve shape is rounded in the elastic-plastic transition region; see figure 1 [4]. In the figure, curves from heat treated sand cast samples (aged at different times at 175 °C) are shown. In the same figure a curve from a 5 mm HPDC specimen is also shown for comparison. It can be observed that the 5 mm HPDC specimen differs from the other curves in terms of curve shape. Another clear example of this difference is seen in figure 2, showing some flow curves for AM-alloys at different Al-contents, published by Hydro Magnesium [5]. It is clear that these curves have a much more defined bend when the material starts to plastify. The reason for this very significant difference in initial flow behaviour has (to the knowledge of the authors) not been discussed in the literature. A tentative explanation for the differences could be related to mould constraints, which will give rise to some plastic deformation during cooling of the tensile test bars in the die.

502 If this plastic strain exceeds only about 0.2-0.5%, it is natural that the tensile test will give rise to a curve with a more significant bend than if no plastic deformation had taken place. If this is the case, it should be realised that published data for the yield strength evaluated from as-cast bars in metal dies is not necessarily representative for castings unconstrained during solidification. Consequently, an un-constrained casting may have lower yield strength than the published value, even though the microstructure is identical. In the present study, some initial experiments were performed to test this hypothesis. 400

FQ2

true stress (MPa)

300

42 h

120 h

16 h

4h T4

hpdc 5mm

200 ac

100

0 0.00

0.04 0.08 true plastic strain

0.12

Figure 1: Examples of true stress – true plastic strain flow curves of AZ91 alloy, as cast (ac), as quenched (T4) and aged for different times (h) at o 165 C. A flow curve of a die cast specimen (hpdc5mm) is included for comparison [4].

2

Figure 2: Stress-strain curves of AM alloys from HPDC as-cast tensile test bars [5].

Experimental procedure

In this study, magnesium alloy AM50 rods having a diameter of 10 mm and a length of 180 mm were produced by re-melting and solidification in a gradient furnace. The advantage of using this procedure is that the scale of the microstructure can be easily controlled simply by varying the drawing rate in the furnace. Additionally, a very low defect level is obtained, thereby resulting in reproducible samples with a very low scatter in results. The cast rods were then machined to the dimensions shown in figure 3. A more detailed description of the experimental procedure has been published elsewhere [2]. Tensile testing bars were also machined from a commercial automotive component of alloy AM50 produced by pressure die casting; see figure 4. Finally, some of the produced tensile testing bars were heated in a furnace to 410 °C for 10 minutes, and then fixed in a vice during cooling in order to simulate the mould constraint which exists in as cast tensile test bars in a metal die. All tensile testing was carried out on a Lloyd EZ50 tensile testing machine at a crosshead speed of 0.5 mm/min using an attached extensometer.

503

Figure 3: Dimensions of machined tensile test bar

Figure 4: Automotive casting, showing position from where the test bars were machined

3

Results

3.1

Tensile test bars from a commercial casting

It is relevant to study if the curve obtained in tensile testing a bar from a HPDC commercial casting has a similar shape to those shown in figure 2. In figure 5, the stress-strain curve from a test bar machined from the casting shown in figure 4 is shown (solid line), together with the curve from a test bar produced in the gradient furnace at a drawing rate of 6 mm/s (dashed line), which resulted in a microstructure similar to that of the casting. It is seen that they correspond well, and that the curve from the component does not have the well-defined bend as was obtained for the HPDC bars tested. Of course, it seems that there is also some kind of mould constraint in the automotive component (see figure 4). However, it is considerably lower than the extreme conditions in tensile test bar geometry.

Figure 5: Stress-strain curves for AM50 from commercial HPDC casting (solid line) and from sample produced in gradient furnace (dashed line)

3.2

Figure 6: Stress-strain curves for AM50 from samples produced in gradient furnace, and exposed to different degrees of pre-strain before tensile testing

Pre-straining of test bars in tensile testing machine

The simplest way to investigate the stress-strain curve shape after some degree of plastic deformation, such as that which could take place during cooling in the die, is of course to pre-strain the test bar and then unload it before the subsequent tensile testing. Test bars

504 were produced in the gradient furnace at a drawing rate of 6 mm/s. Three bars were prestrained to different deformations; 0.5 mm, 0.75 mm and 1 mm, and then drawn until failure. The resulting curves are shown in figure 6 together with a curve from a bar drawn without any pre-strain. The pre-strained curves have been displaced in the X-direction according to their respective plastic pre-strain. As expected, the pre-strained curves have a well-defined bend and after yielding they all fall on the curve for the non-pre-strained sample. Observe the striking similarity with the curves in figure 2. 3.3

Heating and subsequent fixing of tensile test bars during cooling

Another way to simulate what happens due to mould constraint during cooling in a metal die is to heat the samples and thereafter prevent the free shrinkage by fixing them in a vice after removal from the furnace. A drawback of using this method comes from the heat treatment of the microstructure which takes place during heating, and which would make any comparison with the original material impossible. This problem was avoided by heating two identical tensile test bars (gradient furnace, 6 mm/s) for 10 minutes at 410 °C, and then allowing one of the bars to contract freely, while the other was constrained during cooling. However, due to the small dimensions of the bars, the air cooling can be assumed to be rather quick, and therefore the temperature in the bar when it was fully fixed was probably considerably lower than 410 °C. The exact temperatures could not be recorded due to experimental difficulties. From the tensile testing curves of the “heated” and “heated and fixed” samples presented in figure 7, it becomes clear that the mould constraint has a significant influence on the yield strength of the material. The “heated and fixed” sample also has a significant bend in accordance with the curves from the HPDC tensile test bars in figure 2. A very rough estimate of the plastic strains induced during cooling, assuming full mould constraint, can be done by simply multiplying the coefficient of thermal expansion with the assumed cooling interval. Taking the initial temperature as 350 °C and the final temperature as 25 °C (room temperature) gives (300 – 25) * 26.1e-6 = 0.72 %, which is certainly significant, and which indeed should give a sharper yield-point on re-loading.

Figure 7: Stress-strain curves for AM50 from samples produced in gradient furnace. Samples were heated to 410 oC for 10 min prior to cooling with (dashed line) or without (solid line) contraction constraint.

505

4

Discussion

The hypothesis which initiated this work was based on the assumption that die constraint, which gives rise to plastic deformation in the casting during cooling, is the reason for the significant difference in stress-strain curve shape in the elastic-plastic transition region between sand castings and die castings. The initial tests performed all support this hypothesis. It is certainly true that plastic deformation taking place in castings during cooling in constrained moulds is not a newly discovered phenomenon; however, the consequence for the yield strength of the material has not been discussed. Part of the material’s elasticity is of course consumed, accompanied by a deformation hardening process. As a result it is natural that the actual stress when plastic deformation starts in a pre-strained component will be higher than what would be expected from a non pre-strained tensile test bar. Since most data for yield strength of Mg-alloys seem to be evaluated from as-cast HPDC tensile test bars, and that most castings don’t have such a high degree of mould constraint as the test bars, it is very likely that the real yield strength in the casting is lower than expected. From this follows also that the ejection temperature should have an influence on the yield strength; the lower the temperature is, the higher the yield strength as a result of more plastic deformation during cooling in the die. Nowadays, it seems that a majority of design engineers have accepted the fact that the mechanical properties of Mg-alloys vary in cast components as a result of the varying cooling conditions during solidification and subsequent cooling in the solid state. The reason is of course related to the different microstructures produced; i.e. mainly the amount, distribution and fineness of different phases. One complication which also has attracted significant attention during the last years is the fact that optimum mechanical properties, especially the tensile strength and elongation, are very rarely attained in commercial castings due to the existence of different kinds of defects, such as porosity, slag inclusions, segregation bands, hot tears etc. There is still, however, a major lack of knowledge about how much the property values are reduced as a function of size, amount and distribution of these structural defects. Now, considering that any plastic deformation of the casting during cooling in the die will also have an impact on the local yield strength in a component, it follows that any precise prediction of this property will require advanced simulation tools. Today, commercial tools are available which can be used to simulate microstructures and mechanical properties for some cast materials; e.g. MAGMAiron® [6] which for almost 10 years has been used industrially for cast irons. When such a tool becomes available for Mg-alloys it opens up the possibility to link these results with simulations of the stresses and strains produced during the casting process, thereby enabling accurate predictions of how the real yield strength varies in a complex casting. Finally, it should be emphasized that the phenomena discussed above relate not only to Mg-alloys. However, due to the large pre-yield micro strain for Mg as compared to e.g. Al-alloys, which have a more linear elastic behaviour, it is likely that the effect has most importance in Mg-alloy systems.

506

5

Conclusions

This study has gone some way towards understanding the complex interaction between casting process and the local yield strength which can be expected in cast Mg-alloys components. The experiments clearly indicates that any plastic deformation taking place during cooling in the die due to mould constraints, will lead to a higher evaluated yield strength from the stress-strain curve as compared to a non-constrained casting. As a consequence, there is reason to question the validity of published yield strength data for Mgalloys evaluated from HPDC as cast tensile test bars. The questions raised by this study create a need for further investigations.

6

Acknowledgements

The authors wish to express their sincere gratitude to Mr. Olof Granath and Mr. Sven Hedlund at the School of Engineering for invaluable help during the experiments, to Mr. Jan Isaksson and Hans Brantemo at Gjutal AB for valuable discussions and help, and last but not least VINNOVA, The Swedish Agency for Innovation Systems, for financially sponsoring the project.

7

References

[1] W.P. Sequeira, G.L. Dunlop, M.T. Murray, Proc. 3rd Intl. Magnesium Conf, Institute of Materials (UK), 1997, 63-73. [2] H. Cao, M. Wessén, Metallurgical and Materials Transactions, accepted 2003. [3] C. H. Cáceres, C J Davidson, J. R. Griffiths, R. Svensson, Proc. 33rd Australian Foundry Institute Convention, Noosa, Queensland, Oct. 2002. [4] C. H. Cáceres, J. R. Griffiths, C. J. Davidson, C. L. Newton, Mater. Sci. Eng. A, 2002, 325, 344-355. [5] T. K. Aune, H. Westengren and T. Ruden, Proc. Magnesium Properties and Applications for Automobiles, Society of Automotive Engineers, Inc. (USA), Technical Paper #930418, 1993, 51-57. [6] MAGMA GmbH, Aachen, Germany.

507

Understanding the Corrosion Mechanism: a Framework for Improving the Performance of Magnesium Alloys Guangling Song and Andrej Atrens CRC for Cast Metals Manufacturing (CAST), Division of Materials, The University of Queensland, Brisbane Qld Aus 4072

Abstract This paper provides a succinct but nevertheless complete mechanistic overview of magnesium corrosion. This provides an understanding of the types of corrosion exhibited by magnesium alloys, and also of the environmental factors of most importance. This deep understanding is required as a foundation if we are to produce more corrosion resistant alloys. This present analysis provides a theoretical framework for further research. There is still vast scope both for better understanding of corrosion processes as well as on the engineering usage of magnesium. We live in interesting times regarding the study of magnesium corrosion and expect to see significant advances in the next decade!

1

Introduction

Magnesium is a reactive metal and corrosion protection is of great importance [1, 2, 3] particularly for its greatest single market, the automobile industry [4] where there has been a 20% annual growth. Song and Atrens [1] and Ghali [2] have reviewed the corrosion of magnesium. This paper builds on the prior reviews and provides a succinct overview of the corrosion mechanism based on [5, 6, 7, 8, 9, 10, 11, 12, 13, 14, 15] and the literature. We provide a theoretical framework to facilitate understanding of the various corrosion processes. We start by outlining the key features of the corrosion mechanism. Then we outline how these key features have implications for all types of corrosion experienced by magnesium alloys. This mechanistic framework provides a basis for the understanding the following manifestations of magnesium corrosion. • The negative difference effect and hydrogen evolution. • General corrosion and localised corrosion. • Alloying influences on corrosion performance. • Possibilities for the corrosion performance of bulk amorphous magnesium alloys. • Measurement of the corrosion rate by the amount of hydrogen evolved. • Severe problems with the electrochemical methods for corrosion measurement. • Atmospheric corrosion. • Macro-galvanic corrosion when magnesium is coupled with a metal like steel. • Micro-galvanic corrosion associated with second phases. • Environment influences on the corrosion performance. • Stress corrosion cracking and corrosion fatigue.

508 This deep understanding is required to produce magnesium alloys much more resistant to corrosion than the present alloys. We provide a theoretical framework to inform and guide current and future research. This paper is a cut down version, the fuller version is to appear in Advanced Engineering Materials.

2

Corrosion Mechanism

The corrosion mechanism involves all physical and chemical features and processes when magnesium is exposed to an environment, as illustrated in Fig. 1. (a) At a potential more negative than Ep, a protective film covers the surface. (b) At a potential above Ep, the film is partially protective. Corrosion, at breaks in the surface film, occurs by means of the electrochemical partial reactions (1) and (2) and the chemical reaction (3). The product formation reaction, equation (5) repairs the film. (c) The chance development of localised corrosion leads to (d) & (e) the undermining and falling out of particles, even for the corrosion of pure magnesium. 2.1

Key points

i. A partially protective film covers the surface of magnesium. ii. Hydrogen evolution is intimately associated with magnesium dissolution in two separate ways. (a) An electrochemical reaction governed by equation (1) balances the magnesium dissolution reaction, equation (2). (b) Hydrogen is also produced directly in the reaction + of Mg with water, equation (3). Note that the overall reaction, equation (4) produces one molecule of hydrogen gas for each atom of magnesium dissolved. Furthermore, the overall reaction consumes H+ and produces OH-, i.e. the pH increases, which favours the formation of a magnesium hydroxide film by the precipitation reaction, equation (5). 2H+ + 2e = H2 2Mg = 2Mg+ + e 2Mg+ + 2H2O = 2Mg++ + 2OH- + H2 2Mg + 2H+ + 2H2O = 2Mg++ + 2OH- + 2H2 Mg2+ + 2OH- = Mg(OH)2

(cathodic partial reaction) (anodic partial reaction) (chemical reaction) (overall reaction) (product formation)

(1) (2) (3) (4) (5)

iii. Magnesium has a negative free corrosion potential, Ecorr, with a slightly more negative pitting potential, Ep, in solutions of practical importance like 5% NaCl. iv. The chance development of areas of localised corrosion leads to undermining and falling out of particles of magnesium, even for the corrosion of pure magnesium. 2.2

Implications

A. The negative difference effect (NDE) = i + ii + iv. This is shorthand for saying that NDE arrises from an interaction of key aspects i, ii and iv The partially protective film is potential dependent, in that there is essentially a complete film coverage over the whole surface and a low rate of corrosion for potentials below the pitting potential, at the pitting potential there are some film free areas, and the surface area free of surface film increases with increasing potential. As the potential increases above the pitting potential, there is an increasing surface area free of film on which reactions (1) and (2) can occur more freely than on the surface covered by a surface film. Furthermore, as

509 the potential increases above the pitting potential, the speed of reaction (3) is increased and concomitantly, there is an increase in the amount of hydrogen evolved. These two effects, together with particle undermining are the key elements of the negative difference effect. This effect is explained in more detail in the relevant section in the description of the corrosion of pure magnesium.

Mg+ H2

H2 O

Mg2+ + H2

RDS

H2 Mg

Mg

E<Ep

E≥ Ep

(a)

(b)

Mg

E ≥ Ep early stage

Mg

E ≥ Ep late stage or E>>Ep (d) Fig. 1 The corrosion mechanism.

Particle undermining (c)

Mg

E ≥ Ep late stage or E>>Ep (e)

510 B. Localised corrosion = i + iv. This is shorthand for saying that localised corrosion arrises from an interaction of key aspects i and iv. The existence of a partially protective film, and the way chance determines the progression of corrosion means that “general corrosion” is not an issue. The common form of corrosion is localised corrosion, which for magnesium is much different to that of stainless steels C. Impurity tolerance limits = i + ii + iii. The lack of a surface film on the impurity particles together with the negative corrosion potential allows them to be efficient cathodes for hydrogen discharge, thereby providing significant micro-galvanic acceleration of corrosion. This is one forms of micro-galvanic corrosion. D. Influence of Al, Zr, Y… in improving corrosion performance = i + C. Improvement in the corrosion performance of single phase alloys can be expected from an improvement in the surface film if combined with removed or passivated impurities. E. Bulk amorphous alloys and better corrosion performance = i + C + D. A step improvement of corrosion performance is expected from a combination of a more uniform protective film, a higher tolerance limit and a better alloying effect. F. Corrosion rate measurement = ii + A. Hydrogen evolution as an integral part of reaction (4) provides an elegant method by which to measure the corrosion rate of magnesium alloys. The existence of the NDE means that there are severe problems with the electrochemical methods for corrosion measurement. G. Atmospheric corrosion = i + ii. The overall reaction, equation (4) tends to increase the local pH at cathodic sites. This tendency is particularly strong in a thin surface water layer as is often associated with atmospheric corrosion, in which it is relatively easy for reaction (4) to produce a high local pH value, which tends to facilitate the precipitation of a corrosion product film. Consequently, the surface film provides good protection in many cases of atmospheric exposure. H. Macro-galvanic corrosion = iii. Macro-galvanic corrosion occurs when magnesium is coupled with a metal like steel because Mg is the most active engineering metal and consequently the corrosion potential is more negative (i.e. more active) than all the other engineering metals. I. Micro-galvanic corrosion = iii. The anode matrix has a corrosion potential lower than that of the second phases, because each second phase is formed by the reaction of magnesium with a less reactive metal. J. Stress corrosion cracking (SCC) and corrosion fatigue (CF) = i + ii + iv. Material properties are changed by hydrogen entry into the magnesium matrix through a partially protective film. The localised corrosion associated with the chance development of corrosion leads to stress concentrations, and provides easy points for hydrogen entry in film free areas.

3

Corrosion of Pure Magnesium

Magnesium dissolution in aqueous environments proceeds by an electrochemical reaction with water to produce Mg(OH)2 and H2. Magnesium has a standard electrode potential of 2.37 Vnhe. The corrosion potential is ~ -1.7 Vnhe. Magnesium forms a partially protective magnesium hydroxide film. Thermodynamics shows that Mg++ is the stable species in most aqueous solutions up to a pH of about 10, above which Mg(OH)2 is the stable species. Magnesium has a strange phenomenon, the negative difference effect (NDE) for which Song et al [6] proposed a new mechanism, Fig. 1(b).

511

4

Localised Corrosion

For magnesium, corrosion typically takes the form of localised corrosion [5, 6] as indicted by the corrosion mechanism, Fig. 1. On the surface, there is a partially protective film. The corrosion potential is more positive than the pitting potential, for both single phase alloys as well as for two phase alloys. The localised corrosion initiates as irregular pits, which spread laterally and cover the whole surface. There is not much tendency for deep pitting. The reason is that the cathodic reaction, equation (1), produces OH- ions (or equivalently the consumption of H+ ions) with an increase of the pH, stablisation of the magnesium hydroxide film and a decrease in corrosion tendency. Thus localised corrosion in magnesium has an inherent tendency to be self-limiting. This is in marked contrast to the stainless steels, where the occluded pit cell becomes more aggressive and accelerates pitting. However, unlike stainless steels, the localised corrosion in magnesium spreads laterally. It is common for such corrosion to link up and undermine particles, which fall out of the surface, Fig. 1. This particle undermining is common even in the corrosion of pure magnesium.

5

Tolerance Limits

The influence of the contaminants (Fe, Ni, Cu & Co) is also explained by the corrosion mechanism, Fig. 1. The cathodic partial reaction, hydrogen evolution, is particularly important because of the negative corrosion potential. For each of these elements a tolerance limit can be defined. When the impurity content exceeds the tolerance limit, the corrosion rate is high, otherwise the corrosion rate is low. The various mechanistic possibilities are as follows. In each case, a key part of the corrosion acceleration is easy hydrogen evolution on a new phase in contact with the magnesium matrix. • accelerated galvanic corrosion of the magnesium alloy matrix due to general dissolution of the magnesium alloy followed by the re-precipitation of metallic Fe, Ni & Cu on the alloy surface. This mechanism is similar to the acceleration of corrosion by alloyed copper for aluminium alloys. • galvanic acceleration of corrosion when any of the impurity elements Fe, Ni or Cu exceeds its solubility limit and precipitates as a separate phase, δ. • the δ phase is an active cathode for some compositions, but is passive for others. For example, addition of Mn might cause the change from δ to δ’, such that δ’ has electrochemical parameters that make hydrogen evolution difficult on its surface. (Critical parameters to determine would be Ecorr, io and βc).

6

Alloying for Corrosion Resistance

The corrosion mechanism implies that there are two important factors for the influence of alloying: films and hydrogen evolution. Alloying elements that have been associated with increased corrosion resistance include (1) Al and (2) Zr in ingot alloys and (3) passivating elements like Ni and the rare earth elements (RE) Y and Nd in amorphous alloys. Any beneficial influence of these alloying elements has been attributed to improving the protective properties of the surface film. It is particularly interesting that Ni may be beneficial in large concentrations in an amorphous alloy whereas there is the major detrimental influence of Ni in ingot alloys.

512 6.1

Ingot Alloys

The influence of aluminium is difficult to disentangle from the influence of increasing amounts of the second phase. Data [16] for the alloys of the ASxx, AMxx and AExx alloy systems shows a decrease in the corrosion rate with increasing Al content. However, these alloys are two phase and for aluminium contents above 2% there is an increase in the amount of the beta phase. Song et al [12] studied Mg-Al alloys heat treated to be single phase and also found a beneficial influence of aluminium. The mechanism by which Zr improves corrosion resistance was investigated by Song and StJohn [14] using two versions of the alloy MEZ, containing Zr grain refiner (MEZR) and without Zr additions (MEZU). MEZR contained 0.004% Fe and 0.6% Zr whereas MEZU contained 0.013% Fe and 0.005% Zr. This illustrates that one influence of Zr was to react with Fe in the melt, which removed Fe from the alloy and produced a “high purity” alloy. This was part of the explanation for the improvement in the corrosion performance. There was also an additional effect of the second phase, which was more corrosion resistant than the matrix phase. The second phase contained substantial amounts of Zr and RE elements, and the increase in corrosion performance was attributed partly to a more protective film. 6.2

Amorphous Alloys

The influence of passivating elements has been studied by [17] using alloys produced by the rapid solidification (RS). This work showed that Mg65Cu25Y10, Mg65Ni20Nd15 and Mg82Ni18 exhibited passivity over a considerable potential region. However, they were not passive in more concentrated solutions of more practical importance. Their bulk corrosion behaviour showed a similar trend, from Table 1. Table 1. Corrosion rates of selected Mg alloys.

Alloy High purity Mg Surface of die-cast AZ91D β-phase Mg17Al12 RS Mg65Ni20Nd15 RS Mg79Cu21 RS Mg65Cu25Y10 RS Mg67Ni18Nd15 6.3

Environment 1M NaCl 1M NaCl 1M NaCl 0.01M NaCl pH12 0.01M NaCl pH12 0.01M NaCl pH12 3% NaCl pH12

CR, μm/y 1,140 660 300 2.5 410 250 750

Ref [1] [1] [1] [17] [18] [18] [19]

Strategy to Improve the Corrosion Behaviour of Mg Alloys

A step improvement is to be expected for a Mg-X (or Mg-A-B-X) alloy where (1) the alloying element(s) is (are) in solid solution, and (b) the alloy forms a stable passive film based on an oxide of X rather than the Mg-based oxide/hydroxide usually formed on Mg alloys [7]. This is the approach in Fe-Cr alloys, the basis of stainless steels. In Fe-Cr alloys, the nature of the film changes from one based on Fe to a film based on Cr (see eg [20, 21]). This change in film characteristics occurs at about 12% Cr, and concomitantly the corrosion resistance increases by orders of magnitude. The corrosion behaviour is similar to that of Cr, even though there may be only 12% Cr in the alloy. It is suggested that the most exciting passivating elements to study for Mg-base alloys are Cr, Ti and Si. Cr is well known for its passivating properties in highly corrosion resis-

513 tant stainless steels (Fe-Cr alloys). Similarly Ti forms corrosion resistant Ti alloys and Si is important in imparting high corrosion resistance in high Si cast irons. The alternatives Ta, Zr, Hf, Al or Sc are not considered because of various obvious reasons. Production of Mg alloys of macroscopic size containing significant amounts of Cr, Ti or Si in complete homogeneous solid solution has become possible using bulk-amorphous alloys. As it is easy to produce Mg65Cu25Y10 as an amorphous alloy, it is probable that it is possible to produce amorphous versions of Mg65Cr25Y10, Mg65Cr25Ca10, Mg65Ti25Ca10 or Mg55Li20B10Si10C5. Such alloys are at the centre of this suggestion but have not been produced or studied. Such alloys enable the study of the influence of Cr, Ti and Si. A twostep approach is suggested. Step one should explore the properties of such alloys. If indeed there is a step increase in corrosion performance (as might be expected in analogy with the stainless steels) then step two. Step two is the optimisation of other properties, for example decreasing the density, mechanical properties, glass forming ability, processing etc.

7

Measurements of the Corrosion Rate

The corrosion mechanism, Fig. 1, has implications for the measurement of corrosion of magnesium alloys. In particular, (1) there are severe problems with the electrochemical methods for measurement of the corrosion rate, and (2) there is a particularly elegant and easy manner to measure the corrosion rate of magnesium by means of the amount of hydrogen evolved [11]. The principle is based on the overall corrosion reaction, equation (4), which indicates that the dissolution of one atom of magnesium generates one molecule of hydrogen. Therefore, the measurement of the hydrogen evolved is equivalent to the measurement of the weight loss of metal, provided both are converted into the same units. Reaction (4) indicates that the hydrogen evolution is determined by the dissolution of magnesium. The corrosion products do not influence the relationship. The hydrogen evolution directly reflects the corrosion rate of magnesium. Therefore, this method to measure the corrosion rate should be reliable, easy to use and free from many of the errors inherent in the measurement of corrosion by weight loss. A simple way in estimating the error for magnesium alloys is to take the simplest view and assume that the alloy dissolves homogeneously. As all magnesium alloys have small amounts of alloying, less than 10%, the error must be 10% or less, on the basis that the alloying elements do not themselves corrode with the concomitant liberation of hydrogen. Since the corrosion reaction for many of the alloying elements, like Al, is analogous to that for magnesium, the error will in practice be much less Comparison of the corrosion rates of various specimens measured by weight loss and by the hydrogen evolution method for various magnesium alloys, including pure magnesium, single-phase alloys and various forms of AZ91, showed that there was satisfactory agreement between the two methods [11]. This method also has the advantage that the change in corrosion rate can be monitored. This allows the study of how the corrosion rate changes with such variables as time of exposure and changing exposure conditions. The problems associated with the electrochemical methods of measuremen was illustrated by [11], which presents polarisation curves for MEZ and AZ91D in 5% NaCl (pH11). Tafel extrapolation (for both alloys of both the anodic and cathodic polarisation curves) would indicate almost identical corrosion rates. In contrast weight loss measurements indicated that the ratio of weight loss for these two alloys was MEZ:AZ91D = 25:1.

514 The reason that the Tafel extrapolations did not give correct estimates was that the corrosion of magnesium violates the requirements of the Tafel extrapolation method. The same reasons also apply to the use of polarisation resistance.

8

Atmospheric Corrosion

Overall magnesium alloys have a corrosion resistance that is not as good aluminium alloys, because the surface films on the surface of magnesium alloys are not as stable as those on aluminium alloys. This has restricted the application of unprotected magnesium alloys to benign exposure conditions like indoor (and outdoor) atmospheric exposures. It should nevertheless be emphasised that for these environmental conditions, modern magnesium alloys have corrosion rates lower than those of comparable aluminium alloys, and much lower corrosion rates than carbon steel.

9

Macro-Galvanic Corrosion

The active nature of magnesium means that galvanic effects are an issue. Magnesium is more active than all engineering alloys. Consequently magnesium is the anode and corrodes preferentially in any galvanic couple [15].

10 Micro-Galvanic Corrosion and Second Phases The possible types of behaviour are summarised by the data given in Table 2 for magnesium alloys corroding at their free corrosion potentials in 1N NaCl at pH11. High purity magnesium, taken as the standard, showed a corrosion rate of 1.1 mm/y. A higher corrosion rate was shown by the interior of die cast AZ91D and by the high-purity sand-cast AZ91. The β phase accelerated the corrosion. In contrast, the surface of die cast AZ91D had a corrosion rate lower than that of high purity magnesium. The β phase provided protection as is clear from the still lower corrosion rate shown by pure β. The corrosion rate of the β phase is nevertheless high. For aluminium alloys the corrosion properties are severely degraded by this same phase. This points to the methodology by which to increase the corrosion resistance of magnesium alloys by the alloying to produce a radically different second phase, the Ζ-phase. Much better corrosion becomes possible if this new second phase, the Ζ-phase has a corrosion resistance greater than the β phase. Such a Ζ-phase should be more passive than the β phase and also have a much lower ability to allow hydrogen evolution on its surface. Table 2. Corrosion rates under open circuit conditions in 1N NaCl [18].

Sample LP Mg (240 ppm Fe) HP AZ91 sand cast AZ91D – die cast (interior) High purity Mg AZ91D – die cast (surface) β Phase the Ζ-phase

Rate, mm/y 53 12 5.7 1.1 0.66 0.30 ?? 10-6 ??

Mechanism/comment Impurity accelerated corrosion β accelerates corrosion β accelerates corrosion Standard for comparison β protects Significant corrosion rate Needs to be more passive than the β phase and have a lower ability to liberate hydrogen on its surface

515

11 Environmental Influences The various environmental influences may be understood using the principles of the corrosion mechanism, Fig. 1. Corrosion needs a liquid water phase, so the corrosion rate is low in dry air, dry gasses, and in many organic solvents. Corrosion rates, in aqueous solutions, are low under conditions which promote the formation of a stable corrosion product film: e.g. (1) local conditions of pH above 10.5 (to produce a film of Mg(OH)2), including stagnant pure water, (2) fluoride containing solutions (insoluble MgF2 film), (3) MgSO4 in concentrated sulphuric acid, (4) oxidising conditions free of film breakdown agents like chlorides (eg a boiling solution of chromic acid) can lead to stable oxide/hydroxide films. In contrast, high corrosion rates are experienced under conditions leading to surface film beak-down (e.g chloride solutions). Corrosion rates are also high for magnesium in galvanic contact with heavy metals, or for magnesium with surface contaminated with a heavy metal (eg a magnesium surface shot blasted with iron shot), or magnesium in contact with a solution of heavy metal salts.

12 Stress Corrosion Cracking The corrosion mechanism indicates that stress corrosion cracking (SCC) is a concern. Moreover, there is the expectation that environment assisted cracking (EAC) (including SCC and hydrogen embrittlement (HE)) failures will increase with increased use of Mg alloys in load bearing applications. Magnesium alloys are known to be susceptible to EAC. For example, the authoritative ASM Handbook [22], Vol 13, “Corrosion” on SCC of magnesium alloys can be summarised as follows. Magnesium alloys containing more than 1.5% Al are susceptible to SCC. Wrought alloys appear more susceptible than cast alloys. While there is little documentation of service SCC of castings, laboratory tests can cause SCC at tensile loads less than 50% of yield stress in environments causing negligible corrosion. The low incidence of service SCC failures can be attributable to low actual stresses [22] in service in the past. However, the incidence of EAC is expected to increase because the service conditions for Mg alloys are changing, particularly in the automobile industry where magnesium components are being increasingly used in structural load-bearing applications. Furthermore, increased loading on these magnesium components is a natural progression as designers increase the loads and decrease the section sizes of components as part of the environment imperative to decrease weight. Moreover, the magnesium components in the load bearing applications are increasing in complexity, and this increasing complexity increases the probability of high loads in some parts of such components. Unanswered Issues • Are modern AMxx and Zr containing magnesium casting alloys in practice free of EAC as suggested by the early (uncorroborated) reviews in 1966? • Design allowables for safe service of magnesium alloys, and how these are controlled by the alloy and by mechanical and environment loading conditions. • Mechanical and metallurgical conditions leading to EAC for commercial cast magnesium alloys. How the conditions for EAC (initiation stress, initiation stress intensity factor, crack velocity) depend on (1) alloy chemistry variables (major alloying components such as Al, Zn, Y, RE); and (2) minor additions such as Zr, Zn, Mn, Si, RE in AZxx, AMxx, ASxx, ZExx and WExx alloys, (3) microstructural features such as volume fraction, distribution and electrochemical properties of second phases, particularly continuous eutectic phases, and (4) slip type in the matrix phase, at the isolated second phase particles and at the eutectic.

516

•

Environmental/solution/temperature conditions leading to EAC for commercial cast magnesium alloys.

13 Concluding Remarks This paper has built on the prior reviews and has provided a succinct overview of the corrosion mechanism. This understanding of the corrosion mechanism provides a basis for the understanding the various manifestations of corrosion of magnesium alloys, and provides a foundation on which to develop alloys with much greater corrosion resistance.

14 Acknowledgements The CRC for Cast Metal Manufacturing (CAST) was established under, and is supported in part by the Australian Government’s Cooperative Research Centres Scheme.

15 References [1] G Song and A Atrens, Advanced Engin Materials 1 (1999) 11-33. [2] E Ghali, in Uhlig’s Corrosion Handbook, ed RW Revie, Wiley-Interscience, (2000) p793-830. [3] Y Kojima et al eds, Magnesium Alloys 2003, Materials Science Forum 419-422 (2003)1-1037. [4] GS Cole, Materials Science Forum 419-422 (2003) 43-50. [5] G Song, A Atrens, S StJohn, J Nairn and Y Lang, Corrosion Science 39 (1997) 855. [6] G Song, A Atrens, D StJohn, X Wu and J Nairn, Corrosion Science 39 (1997) 19812004. [7] G Song, A Atrens, X Wu and B Zhang, Corrosion Science 40 (1998) 1769-1791. [8] G Song, A Atrens and M Dargusch, Corrosion Science 41 (1999) 249-273. st [9] G Song, Corrosion Science in the 21 century UMIST 7-11 July, 2003. [10] G Song and DH StJohn, Journal of Light Metals 2 (2002) 1-16. [11] G Song, A Atrens and DH StJohn, in J Hryn ed, Magnesium Technology 2001, New Orleans, TMS (2001) 255-262. [12] G Song, AL Bowles and DH StJohn, Mater Sci and Engineering, accepted for publication [13] G Song and A Atrens, Wolfsburg (1998) 415. [14] G Song and DH StJohn, J Light Metals 2(2002)1. [15] G Song, B Jonhannesson, Sarath Hapugoda, DH St John, Corrosion Science, in press. [16] O Lunder, K Nisanccioglu and RS Hansen, Paper No 930755SAE, Detroit MI, (1993) 117. [17] HB Yao, Y Li and ATS Wee, Electrochim Acta 48 (2003) 2641 [18] HB Yao, Y Li, ATS Wee, JS Pan and JW Chai, Surface Review Letters 8 (2001) 575. [19] HB Yao, Y Li, ATS Wee, JS Pan and JW Chai, Applied Surface Science 158 (2000) 112. [20] P Bruesh, K Miller, A Atrens and H Neff, Appl Phys A 38 (1985) 1-18. [21] S Jin and A Atrens, Applied Physics A 42 (1987) 149-65. [22] ASM Handbook, Vol 13, Corrosion ASM International, fourth printing (1992).

517

Potentiodynamic Studies of Some AZ and ZA Magnesium Alloys in Different Corrosive Media Mustafa Özgür Öteyaka, Anne-Marie Lafront, Edward Ghali and Réal Tremblay Université Laval, Québec, Canada

1

Introduction

It is well documented that the presence of heavy metallic impurities like Fe, Ni and Cu represents the most detrimental factor influencing the corrosion properties of commercial magnesium alloys [1-3]. The corrosion behaviour of cast Mg-Al alloys could depend considerably on the microstructure (presence of β-phase and impurities) and the environment to which they are exposed. Earlier investigation on the corrosion behaviour of AZ91 alloys proved that the β-phase could play a dual role in the dissolution behaviour since it can act as either a galvanic cathode or a kinetic barrier to dissolution [4-7]. Changes in aluminum and zinc contents and/or addition of other elements in magnesium alloys tend to modify significantly the microstructure and influence notably the corrosion resistance in specific media. In this work, electrochemical techniques were used to compare the corrosion behaviour of two groups of magnesium alloys: 1/AZ91D and AZ91E, and 2/ZA104, ZAC10403 (0,3% Ca), and ZACS1040305 (0,3% Ca and 0,5% Sr) in diverse conditions and media.

2

Experimental Procedure

2.1

Material used and specimen preparation

Specimens were prepared from cast parts of two groups of magnesium alloys. The AZ group had aluminum as principal alloying element (~9,0% Al) whereas the ZA group had zinc (~10,0% Zn). The AZ91D, AZ91E, ZA104, ZAC10403 and ZACS1040305 alloys were cast in permanent mould in accordance to the standard ASTM B93 [8]. The chemical composition of each alloy is given in Table 1. Each corrosion specimen consisted of 1 cm x 1 cm x 2 cm metallic piece which was encased in epoxy resin. The surface analyzed was polished with 600 and 1200 grit SiC abrasive papers, cleaned with alcohol and then dried before immersion in the test solution. Table 1. Chemical composition (in wt%) of the experimental magnesium alloys.

Alloys AZ91D AZ91E ZA104 ZAC10403 ZACS1040305

Al Zn 9,00 0,67 8,70 0,70 3,69 9,73 3,78 10,50 4,12 11,70

Ca ------0,31 0,31

Sr --------0,54

Mn 0,33 0,26 0,25 0,03 0,35

Si 0,010 0,200 0,015 0,001 18 ppm

Fe 0,005 0,005 0,030 0,003 189 ppm

Cu 0,030 0,015 < 0,001 < 0,001 2 ppm

Ni 0,002 0,001 0,001 0,002 2 ppm

Mg bal. bal. bal. bal. bal.

518 2.2

Testing corrosive solutions

Three different corrosive solutions were used for this study with or without agitation: 5% NaCl solution saturated with Mg(OH)2 at pH 9, distilled water saturated with Mg(OH)2 solution at pH 9, and 0,01M NaOH solution at pH 12. 2.3

Electrochemical technique parameters

The standard ASTM G5 was used to carry out potentiodynamic tests at 25 °C [9]. The potentiodynamic curves were traced at 0,1 mV/s or 0,07 mV/s scanning rate over a potential range of 1000 mV starting from the cathodic end using a potentiostat/galvanostat (Model 273 EG&G). For some experiments, the solution was stirred at 250 rpm and heated at 75 °C. Also, other scanning rates, potential ranges and cathodic polarization voltages were used. After immersion of the sample into the solution, the test was started within one minute. All potentials were given with respect to a silver chloride reference electrode Ag, AgCl/KClsat. (0,222 V vs. SHE). The reproducibility of the corrosion potential and the corrosion current values were ±10-15 mV and ±10-15% respectively.

3

Results and Discussion

In this paper, the following terms Ecorr., icorr., Ep, ip, Ecp, and icp correspond to corrosion potential, corrosion current, pitting potential, pitting current, critical passivation potential, and critical passivation current respectively. The polarization curve was started from the cathodic region after a stabilization period of one minute in the solution. The corrosion potential was determined from the potentiodynamic polarization curves, using dedicated software. It is arbitrarily admitted that the two parameters Ecp and icp can describe passivation in spite of the observed partial passivation or pseudo-passivation of the magnesium alloys in the examined media. Pitting potential was determined from the potentiodynamic curves at the rapid increase of current density in some potentiodynamic curves. 3.1

Influence of different aqueous media at 25 °C

Results from potentiodynamic studies in the three corrosive media with and without stirring for magnesium alloys are given in Table 2. In NaCl solution, the potentiodynamic tests for ZA104, ZAC10403, and ZACS1040305 alloys exhibited a passivation zone with formation of a black corrosion product acting as a barrier layer followed by dissolution of the surface. But AZ91D and AZ91E alloys did not present a good passivation zone and dissolution of the surface is accompanied by the formation of pits that have been observed microscopically. Active-passive parameters of magnesium alloys in a strong corrosive medium (5% NaCl solution) cannot be determined precisely from potentiodynamic measurements. The values of corrosion potential as determined from the polarization curves show that they are close to each other. However, if the potential of AZ91D is compared to that of ZACS1040305, it can be stated generally and statistically that the alloy containing higher quantities of Zn and small quantities of Ca and Sr have more active potentials. The minimum corrosion current rates are found for the AZ91E followed by AZ91D and somewhat inferior to that of ZAC10403 and ZACS1040305. This is not surprising since AZ91E (with less impurity) has better corrosion resistance than AZ91D. Typical corrosion

519 rates of AZ91E in salt fog have been already reported to be ~0,25 mm/year and this is about a hundred times lower than that of AZ91C standard alloy [2]. Table 2. Data from the potentiodynamic curves obtained at 0,1 mV/s scanning rate and 25 °C in different media without and with (values in parentheses) stirring for magnesium alloys.

icorr. Ep ip Ecorr. (V) (μA/cm2) (V) (μA/cm2) 5% NaCl solution saturated with Mg(OH)2 at pH 9 AZ91D -1,31 (-1,30) 3 (3) ----AZ91E -1,30 (-1,32) 2 (1) ----ZA104 -1,32 (-1,33) --- (3) ----ZAC10403 -1,32 (-1,32) 5 (2) ----ZACS1040305 -1,33 (-1,32) 4 (---) ----Distilled water saturated with Mg(OH)2 solution at pH 9 AZ91E -1,29 (-1,29) 0,1 (0,1) -0,89 (-0,72) 25 (10) ZACS1040305 -1,22 (-1,26) 0,2 (0,2) ----0,01M NaOH solution at pH 12 AZ91E -1,34 (-1,27) 0,1 (0,02) -0,73 (-0,2(1)) 15 (---) ZA104 --- (-1,16) --- (0,03) --- (-0,30) --- ( 5 ) ZAC10403 --- (-1,26) --- (0,03) --- (-0,32) --- (34) ZACS1040305 -1,29 (-1,11) 0,6 (0,01) -0,82 (-0,2(1)) 23 (---) Alloys

(1)

Ecp (V)

icp (μA/cm2)

-----------

-----------

-1,20 (-1,19) 18 (10) -----1,21 (-1,19) --- (-0,99) --- (-0,98) -1,22 (-0,95)

16 ( 2 ) --- ( 5 ) --- (23) 15 (3)

More positive than -0,2 V.

Figure 1 shows the anodic polarization curves associated with some scanned surface views of AZ91E, ZAC10403, and ZACS1040305 alloys. For AZ91E alloy, the pits appear at the beginning of the anodic polarization; they expand and cover large spots on the surface with the increase of the applied current (Figures 1a and 1b "A" and "B"). ZAC10403 and ZACS1040305 alloys exhibit the same corrosion behaviour, at the beginning of anodic polarization, round corrosion products appear on the sample surface (Figures 1a and 1b "D"). However at higher applied current densities, the corrosion products involve their surfaces (Figure 1b "E"). This corresponds to the passive zone on polarization curves which is able to prevent the corrosion rate increase of the alloy.

Figure 1. a) Anodic potentiodynamic curves of the AZ91E, ZAC10403 and ZACS1040305 alloys at scan rate of 0,07 mV/s in the 5% NaCl solution saturated with Mg(OH)2 at pH 9 and 25 °C (w/o stirring); b) surface views at different potential-current points for AZ91E (A and B) and ZAC10403 (D and E) specimens.

520 At 1000 mV potential range, a passive zone and a clear pitting corrosion potential for AZ91E alloy are observed in Mg(OH)2 solution. The alloy ZACS1040305 does not show an active-passive behaviour in Mg(OH)2 solution as seen in NaCl solution. Corrosion rate of ZACS1040305 is two times higher than that of AZ91E. It can be postulated then that the better performance of AZ91E is due to the formation of a better passive layer containing magnesium hydroxide as a barrier that obstructs dissolution [1]. The poor performance of ZACS1040305 alloy in this solution is probably due to the galvanic effect of the intermetallic phases (MgxZnyAlz) with magnesium matrix. The potentiodynamic curves of AZ91E and ZACS1040305 alloys in NaOH solution have shown good passivation zone and an obvious pitting corrosion potential for both alloys. However, it seems that the aptitude to passivation (Ecp - Ecorr.) for AZ91E is more difficult in NaOH solution at pH 12 than in Mg(OH)2 solution at pH 9 although the icp does not give the same trend and the quality of passivation is better. Both alloys have a corrosion potential more negative than that in Mg(OH)2 solution (Table 2). The alloy AZ91E demonstrates a 50 mV more active potential than ZACS1040305, while the corrosion rate of this latter is six times higher than that of AZ91E in NaOH solution. The corrosion rates of AZ91E remains constant in both solutions, while corrosion rate for ZACS1040305 increases three times in NaOH solution as compared to Mg(OH)2 solution. The corrosion rates of AZ91E are similar in the Mg(OH)2 and NaOH solutions but much less than that in NaCl solution. It seems that AZ91E has a better corrosion resistance in Mg(OH)2 because of its aluminum content (9%). The ZACS1040305 experimental alloy has more noble corrosion potential and less corrosion rate in Mg(OH)2 than in NaOH and NaCl solutions. This may indicate that ZACS1040305 corrosion resistance is increased more by the existence of the Mg(OH)2 at the interface than by a more alkaline pH [4,10-12]. 3.2

Influence of agitation

Corrosion behaviour is observed for the two groups of magnesium alloys in non-stirred and stirred NaCl solutions (Figures 1 and 2). In general, the corrosion potentials had respectively similar values (Table 2).

Figure 2. Potentiodynamic curves for magnesium alloys at scan rate 0,1 mV/s in 5% NaCl saturated with Mg(OH)2 stirred solution at pH 9 and 25 °C.

521 For AZ91E alloy, stirring shows an improving effect of the passivation zone in Mg(OH)2 solution. Also it is found that the corrosion potential is equal in stirred and nonstirred solutions. Similar corrosion rate values are found in the stirred and non-stirred solutions for both alloys (Table 2). For ZACS1040305 alloy, a weak passivation zone is observed in stirred solution. For the two groups of alloys, agitation shows a favorable influence on the passivation zone in NaOH solution. The inferior passive domain of ZAC10403 (660 mV) compared to ZA104 (690 mV) explains the no good effect of 0,3% Ca on passivation. The strontium associated with calcium has more evident influence on passivation than calcium alone. The alloy AZ91E shows a better aptitude to passivation (Ecp - Ecorr.) with a value of 80 mV compared to the group of alloys containing more zinc (Table 2). The corrosion rate (icorr.) of AZ91E and ZACS1040305 alloys in the stirred NaOH solution is generally much less important than that in the same non-stirred solution. In the group of ZA alloys, the corrosion rate of ZA104 and ZAC10403 is three times higher than that of ZACS1040305 in stirred NaOH solution. 3.3

Influence of temperature

In NaCl solution at 75 °C, AZ91E alloy presents an active-passive behaviour that is not observed clearly at 25 °C and dissolution of the surface is accompanied by pitting at both temperatures (Figure 3). The active-passive behaviour at elevated temperature is due probably to the formation of a more stable film on β-phase (Mg17Al12) because of the presence of higher amount of aluminum. The continuous and stable oxide film acts as barrier so that the dissolution of α-phase (α-Mg) is inhibited. At 25 °C, the ZACS1040305 alloy shows a relatively better active-passive behaviour than that at 75 °C. Increasing the temperature drops the corrosion potential of the alloy AZ91E towards more negative values (Δ = 40 mV). However, the corrosion potential of the ZACS1040305 alloy shifts in the opposite direction and becomes 30 mV nobler. At 75 °C, the corrosion potential of the alloy ZACS1040305 is 40 mV nobler than AZ91E alloy. The corrosion rate of the alloy ZACS1040305 at elevated temperature is half of that at ambient temperature (Table 3). Table 3. Data from the potentiody-namic curves presented in Figure 3.

Figure 3. Potentiodynamic curves for different alloys at scan rate 0,1 mV/s in 5% NaCl solution saturated with Mg(OH)2 at pH 9 and 25 °C - 75 °C (w/o stirring).

Alloys

Ecorr. (V)

icorr. (μA/cm2)

AZ91E (25 °C) AZ91E (75 °C) ZACS1040305 (25 °C) ZACS1040305 (75 °C)

-1,30 -1,34 -1,33

2 3 4

-1,30

2

522 It can be postulated that increasing temperature for the AZ group of alloys shifts the corrosion potential towards more negative values, increases the corrosion rate, and favours the active-passive behaviour. An opposite tendency is observed for the ZA alloys for which a shift of the corrosion potential to more noble values, a decrease of the corrosion rate and disappearance of the active-passive behaviour appear with an increase of the temperature. 3.4

Influence of scan rate

ZACS1040305 alloy was tested at a cathodic range of -40/250 mV for each scan in NaCl solution at 25 °C without stirring. The passivation zone appears more visible with a low scan rate but it is less evident at 0,5 mV/s (Figure 4). At low scan rate of 0,01 mV/s, the best aptitude to passivation is observed. Also the critical passivation current cited in Table 4, shows that the aptitude to passivation is the best at 0,01 mV/s (4100 μA/cm2). It is recommended to choose a scan rate between 0,01 and 0,1 mV/s for the alloy ZACS1040305 to perform the potentiodynamic tests, although the polarization studies for AZ91E alloy show very similar curves at the three scan rates used.

Figure 4. Potentiodynamic curves for ZACS1040305 alloy at different scan rates in 5% NaCl solution saturated with Mg(OH)2 at pH 9 and 25 °C (w/o stirring). Table 4. Data from the potentiodynamic curves presented in Figure 4.

Scan rate (mV/s)

Ecorr. (V)

icorr. (μA/cm2)

Ecp (V)

icp (μA/cm2)

0,01 0,1 0,5

-1,30 -1,34 -1,30

--1,3 ---

-1,30 -1,29 ---

4100 15000 ---

4

Conclusions

The electrochemical studies prove that, without agitation, the corrosion rates of AZ91E are similar in distilled water saturated with Mg(OH)2 at pH 9 and 0,01M NaOH at pH 12 solutions but much less than that in 5% NaCl solution saturated with Mg(OH)2 at pH 9.

523 The ZACS1040305 experimental alloy has more noble corrosion potential and less corrosion rate in Mg(OH)2 solution than in NaOH solution and 5% NaCl solution saturated with Mg(OH)2 . For both groups of alloys, stirring seems to have no effect on corrosion behaviour in 5% NaCl solution saturated with Mg(OH)2, but helps to increase the passive domain of the alloy AZ91E in the Mg(OH)2 solution. Generally for AZ91E and ZACS1040305 alloys in NaOH solution, agitation decreases the corrosion rate and critical passivation current and increases the passive domain. For the AZ alloys, increasing temperature shifts the corrosion potential towards more negative values, increases the corrosion rate, and favors the active-passive behaviour in 5% NaCl solution saturated with Mg(OH)2. However, the opposite tendency is observed for the ZA alloys. It is also recommended to choose a scan rate between 0,03 and 0,1 mV/s for the alloy ZACS1040305 to perform the potentiodynamic tests in 5% NaCl solution saturated with Mg(OH)2. For the studied compositions of ZA alloys, it should be noted that the addition of Sr associated with Ca improves the corrosion resistance better than the addition of Ca alone. Finally, increasing the Zn content and lowering the Al content in ZA alloys relatively to the AZ group do not seem to improve the corrosion resistance.

5

Acknowledgement

This work was supported by the Natural Sciences and Engineering Research Council of Canada (NSERC).

6 [1]

References

M. M. Avedesian, H. Baker, Magnesium and Magnesium Alloys, ASM Specialty Handbook, OH, USA, 1999. [2] K. Nisancioglu, O. Lunder, T. K. Aune, Inter. Magnesium Ass., 1990, 43-50. [3] J. D. Hanawalt, C. E. Nelson, J. A. Peloubet, Trans. AIME, 1942, 147, 273-299. [4] H. Alves, U. Koster, E. Aghion, D. Eliezer, Materials Technology, 2001, 16, 2, 110-126. [5] R. Ambat, N. N. Aung, W. Zhou, Corrosion Science, 2000, 42, 1433-1455. [6] G. Song, A. Atrens, M. Dargusch, Corrosion Science, 1999, 41, 249-273. [7] O. Lunder, J. E. Lein, T. Kr.. Aune, K. Nisancioglu, Corrosion, 1989, 45, 741. [8] ASTM Standard B93/B93M-94b, Annual Book of ASTM Standards, 02.02, 1997, 46-49. [9] ASTM Standard G5-94, Annual Book of ASTM Standards, 03.02, 1997, 54-64. nd [10] P. Uzan, N. Frumin, D. Eliezer, E. Aghion, Proc. 2 Israeli Inter. Conf. on Magnesium Science & Technology, MRI, Beer-Sheva, Israel, 2000, 385-391. [11] O. Lunder, M. Videm, K. Nisancioglu, Corrosion Resistant Magnesium Alloy, SAE Technical Paper 950428, Detroit, USA, 1995, 57-62. [12] O. Lunder, K. Nisancioglu, R. S. Hansen, Corrosion of Die Cast MagnesiumAluminum Alloys, SAE Technical Paper 939755, Detroit, USA, 1993, 117-126.

524

Plating of Magnesium – New Developments Peter Gregg Franz Oberflächentechnik GmbH & Co. KG, Geretsried - Germany

Introduction Magnesium is more and more used in applications, which require high strength of the material combined with low weight. Especially in car industry and electronics industry we can see a boom of using magnesium diecasted parts – mobile phone chassis, complete car doors, gearboxes, internal car parts etc. are now made of Magnesium to reduce weight with no concession regarding the solidity. The special casting properties of Magnesium allow to produce big parts with low wall thickness, so there is much more flexibility regarding the geometrics compared to aluminium diecasting. Due to these casting properties Magnesium can also be used to substitute sheet metal parts in cars (car doors for example).

1

Properties of Magnesium regarding corrosion

One of the main problems with Magnesium parts is corrosion: nearly every part needs some corrosion protection, because the surface of Magnesium is highly reactive and corrodes when getting in contact with water or high humidity. On aluminium parts there is a “natural” oxide layer, which protects the substrate against the environment and gives a good corrosion resistance without any further treatment. On Magnesium parts, most applications need some corrosion protection, because this “natural” protection does not exist. Therefore the surface has to be covered by a coating, which avoids corrosion and meets further demands like abrasive properties, cosmetics, deforming properties etc.

2

Surface treatments of Magnesium

2.1

Electroplating

Process The process of electroplating of magnesium is quite similar to electroplating of Aluminium – after cleaning and activation a thin layer of zincate is put on the part and after that a copper layer of 3 to 10 micron. Onto the copper layer there is a wide range of possibilities for other metal layers or e-coat – depending of the demands and the application you can use nickel, tin, zinc, tin/zinc alloys, copper/tin alloys etc.

525 Properties of surface Depending on the top layer the properties of the surface can be varied in a wide range. If corrosion resistance and conductivity are the major demands, we recommend tin or a tin/zinc alloy, which can be chromated or plated with e-coat to increase corrosion resistance. For decorative applications, nickel plus chrome or Cu/Sn-alloy are the first choice; the properties of this plating technology are quite similar to electroplating of aluminium. The corrosion resistance of this coating can be brought to more that 100 hours salt spray test when using optimized Ni-layers, so it is possible to meet the demands of decorative parts in the automotive industry. Applications Electroplating on Magnesium parts is now mainly used for parts in the mobile phone industry, where cosmetic issues and the conductivity of the surface are most important demands. For decorative applications in the automotive industry (especially parts inside the car) this technology is also already used in serial production for high performance sports cars. Costs The costs of electroplating of magnesium are similar to aluminium plating – if there are high quantities, the plating of magnesium becomes cheaper compared to aluminium, because the process is easier to handle and control. The costs regarding waste water and environment are lower compared to aluminium plating, because the chemicals are not so critical. 2.2

Chromating (black) plus paint/e-coat

Process The process of chromate coating is quite simple: after pre-treatment the surface is chemically converted to a chromate; it is similar to chromating of aluminium or other materials. Properties of surface The chromating layer is quite thin and gives some corrosion protection, but in most cases you need additional coating to meet higher demands. The chromating itself cannot be used for visible surfaces, because it does not look good; also the abrasive properties require additional coating on the chromating. The process and the layer use Cr(6), which can cause cancer. The chromated surface itself gives a corrosion protection of about 10 hours (depending on the casting) in salt spray test. The cromated surface can be coloured matt black, which is quite interesting for applications in the optical industry. Applications Chromating is now used on quite a lot of magnesium parts in the automotive industry as a pre-coat for e-coat or powder-coat, but there is no future for this coating because Cr(6) will be forbidden for automotive applications. In the optical industry there is a market especially for black chromating, because it is easy, cheap and effective.

526 Costs Conversion coating/chromating is cheaper than electroplating; no electricity is needed in the process and the plant is quite easy to install. 2.3

Conversion coating (plus e-coat)

Process The process of conversion coating is similar to chromating, but contains no Cr(6) so it can be used for future applications in the car industry. There are several conversion coatings on the market – there is no significant difference between these coatings but the corrosion resistance of all these conversion coatings is heavily influenced by the pre-treatment: with a perfect pre-treatment it is possible to double the corrosion resistance. Properties of surface The conversion coating layer is quite thin and gives some corrosion protection, but in most cases you need an additional coating to meet higher demands. In most cases the conversion coating is combined with e-coat; but especially for decorative surfaces it is necessary to add a additional process step between conversion coating and e-coat to get a class-A surface, otherwise small defects (i.e. spots) can occur. Applications Conversion coating is now used on quite a lot of magnesium parts in the automotive industry as a pre-coat for e-coat or powder-coat. There are also applications in optical industry or mobile phone industry – in most cases there is a kind of paint as final layer to fulfil cosmetic demands and corrosion resistance. Costs Conversion coating is cheaper than electroplating; no electricity is needed in the process and the plant is quite easy to install. 2.4

Anodising ANOMAG plus paint/e-coat

Process The ANOMAG-layer is build up on the surface (not like anodising of Aluminium) and consists of a mixture of Mg-Oxides and Mg-Phosphates. The natural colour is grey; the surface can be coloured (like anodised Aluminium) and the roughness can be adjusted to get a smoother surface. Properties of surface ANOMAG is a good undercoat for lacquer and adhesives – because of the rough structure of the natural ANOMAG layer the adhesion is very good. Tests show, that the adhesion spliced parts is twice better compared to chromating or conversion coating. The corrosion resistance of ANOMAG plus e-coat shows very good results (1000 h salt pray test) and is better than chromating or conversion coating plus e-coat.

527 Applications The ANOMAG-plating is basically suitable for applications in car industry, aviation industry and for chassis parts in the electronic industry, but until today there is no breakthrough in the market because of the quite high costs compared to conversion coating. Costs The costs for ANOMAG are higher than conversion coating and the handling of the parts on the racks is quite difficult, because the process runs with high voltage and current, so good contact between part and jig is necessary. 2.5

E-coat plus electroplating surface

Process This process combines conversion coating/e-coat and electroplating: the part is first plated with conversion coating and a special e-coat; on top of this e-coat a electroplating layer of Ni plus Cr is build up to achieve cosmetic demands. The electroplating on top of the ecoat can be compared with electroplating of plastics but requires a special combination of e-coat and electroless Ni to get good results. The quality of the e-coat is crucial for the cosmetic demands – it is necessary to have a spot-free surface otherwise the whole process does not give good cosmetic results.

Properties of surface This combination gives very good corrosion resistance because of the e-coat and very good cosmetic properties, because the e-coat also levels the surface. With this process it is possible to produce class-A parts outside the car body with chromium surface. Applications The major applications are automotive parts inside or outside the car – for high performance sports cars there are already parts in serial production. If it is necessary to get a chromium surface on a magnesium part combined with more that 500 h salt spray test there is no other choice today. Costs The costs for this process are quite high, because the parts have to be put into at least 2 different plating lines. Therefore it is only suitable for high-level products with high cosmetic demands. It is possible to reduce these costs by building up a dedicated line for this process, but until now the market is too small for such a production line.

528

3

Evironment and Magnesium Coating

Regarding the environmental issues, the electroplating of magnesium causes the same problems as electroplating of aluminium: toxic chemicals are used and – depending on the required surface – heavy metals like nickel, chromium etc. are in the process and in the surface. The problems during production can be solved by a modern waste water treatment and by training the staff, so that contamination of the environment can be avoided during the process. There are some critical issues of heavy metals in the surface: nickel i.e. can cause allergic reactions, so for future developments we try to find solutions without nickel (i.e. Cu/Sn-alloys). In the chromating process the situation is app. the same because of Cr(6) in the process and the layer. n the ANOMAG-process there are no heavy metals and no chemicals, which cause environmental problems; also the coating itself contains no heavy metals and can quite easy be stripped off by chemicals. Therefore ANOMAG is environment friendly compared to other coatings and gives few problems in the waste water treatment.

4

Summary

In the future Magnesium will be used much more in automotive industry, so the demands for a cheap and corrosion resistant coating of magnesium parts will increase rapidly during the next years. There is still a lot of development work to do to find easy and inexpensive solutions, so that magnesium parts can compete with steel or aluminium. Especially for decorative parts the whole system from casting to final chromium surface needs a lot of fine tuning and coordination to get good results and high yield.

529

Surface Treatment of Magnesium Substrates Michael Walter AHC Oberflächentechnik Holding GmbH, Kerpen

Magnesium-based materials used for structural purposes are used wherever a reduction in mass has priority. With a density of 1.74 g/cm3 they partly show specific strength properties that correspond to those of aluminium or steel, or even higher. Furthermore, magnesium substrates are completely recyclable. Despite the development of high-purity alloys most magnesium-based applications require a surface treatment in order to generally improve the corrosion resistance of this non-precious material used for structural purposes (standard electrode potential of – 2.36 V). AHC Oberflächentechnik, Kerpen, has developed various processes for the surface treatment of magnesium alloys. The two most important processes are introduced in the following. We would like to start with the introduction of a chrome-free passivation process for magnesium substrates [1]. This process produces a conversion layer consisting of oxides from both the base material itself and the oxidic reaction products which derive from components of the aqueous passivation electrolyte. The chrome-free passivation layer serves as an intermediate layer for the subsequent application of paint or adhesive. The purpose of this passivation process developed at AHC Oberflächentechnik is to provide an equivalent alternative to chromate conversion coatings and it can be applied to any magnesium-based material of common use, such as die casting, cast and wrought alloys. The process includes the following steps: degreasing, pickling, passivation, drying as well as either sealing or painting. The passivation is carried out at room temperature and can be applied by dipping, splashing, spraying, painting and other commonly used methods of application. Plant designs as used for chromate conversion coatings may also be used for the new passivation process. The aqueous passivation solution consists of potassium permanganate in combination with the anions vanadate, molybdate or tungstate. By Auger-spectroscopy, MgO, Mn2O3, MnO2 and at least one oxide from the group of vanadium, molybdene or tungsten can be determined in the conversion layer. Thus it is possible to demonstrate that the reaction of the passivation solution is like that of a redox system. Permanganate, for example, is reduced to mangane oxides of low-level oxidation. The reaction mechanism of this system is similar to that of chromate conversion coatings. Within the periodic system of the elements, the used elements of the passivation substances are very close to chromium. Picture 1 shows SEM pictures of a passivated surface on AZ91 HP, treated according to one of AHC Oberflächentechnik’s processes which is ® marketed under the trademark MAGPASS-COAT . These pictures show a finely structured conversion layer, similar to that of magnesium surfaces treated with a chromate conversion coating. The chrome-free passivation process for magnesium substrates of AHC Oberflächentechnik provides an equivalent alternative to chromate for the automo-

530 tive industry, mechanical engineering in general, the optical industry, micro electronics and other industrial sectors. Within the total coating system with subsequent painting, this process helps to provide corrosion protection.

Picture 1: SEM photographs of a AZ 91 HP passivated surface treated with the MAGPASS-COAT® process (Photos: AHC)

The space frame of the Volkswagen 1-Litre car (realized as a design study) is made from magnesium, which is even lighter than aluminium. This magnesium frame (picture 2) has been treated with the chrome-free passivation process developed by AHC Oberflächentechnik. After thorough examination of various chrome-free systems of different suppliers, Volkswagen decided to use the AHC process. The space frame is a composite welding and adhesive bonding construction of different magnesium alloys. The coated surface area amounts to a size of approx. 20 m2. Some areas of the magnesium space frame additionally received a clear paint sealing. About two years ago Audi brought a 12cylinder engine onto the market with its A8 L 6.0 quattro. As well as its driving performance, this limousine is notable for its forwardlooking lightweight construction technology. It is the only 12-cylinder model that makes use of the Audi Space Frame ASF® made from aluminium. Also made of aluminium are the wheel supports, the front and rear axle control arms, all the brake callipers, the shock absorber Picture 2: Cockpit of Volkswagen’s 1-Litre-Car: bearings and the wheels. Consequently, lightThe lower part shows the magnesium space weight metals, namely aluminium and magneframe treated with the chrome-free passivation sium substrates, were used for engine parts. A process MAGPASS-COAT . compact construction combining two 6(Photo: Volkswagen) cylinder V engines to a 12-cylinder aggregate in a W shape, together with the lightweight construction, contributes to the low weight of the engine. An essential contribution to peak automotive engineering performance, but one that is often out of sight, is always surface technology. As an example of one such detail, the inlet pipe of the W 12-cylinder engine may be mentioned here (picture 3). It is made of a high-purity sand casting magnesium alloy – a contribution to weight reduction. On the one hand, to combat the marked corrosion tendency of magnesium and its alloys and, on the ®

531 other hand, to satisfy the requirement for a chrome-free coating, the inlet pipe is treated with an aqueous chrome-free passivation solution. The MAGPASS-COAT® layer alone offers only temporary corrosion protection. In the case of the inlet pipe a 200 μm polyester powder coating in titanium silver colour is applied. This layer combination not only has an attractive appearance but stands up to all the strength and corrosion tests demanded by Audi. Another important AHC process for surface treatment of magnesium substrates, namely ® Picture 3: Inlet pipe W 12 made of AZ 91, MAGOXID-COAT , belongs to the group of treated with the the chrome-free passivation plasma chemical processes. Plasma chemical MAGPASS-COAT and a subsequent polyester coatings are carried out in electrolytes which are powder coating saline solutions. Making use of an external (Photo: AHC) power source the workpiece being processed takes on the function of the anode. Anodizing takes place as the plasma is discharged in the electrolyte on the surface of the workpiece to be treated. The substrate surface is partially melted by bursts of oxygen plasma produced in the electrolyte and an adhesive oxide ceramic metal compound forms on the workpiece. 50% of the crystalline oxide ceramic conversion layer grows into the base material whilst 50% grows outwards. Edges, cavities and relief designs are coated uni® formly. The MAGOXID-COAT layer contains a large proportion of highly resistant elements such as spinels, e.g. MgAl2O4. The produced oxide ceramic layers not only provide protection against wear and corrosion but also fulfil requirements regarding hardness, uniform layer deposition, fatigue strength under reversed bending stresses, dimensional accuracy or temperature load capacities [2 – 5]. MAGOXID-COAT® layers provide a good wear resistance similar to that of AHC’s hard anodizing layers, known on the market as HART-COAT® layers. The corrosion and wear resistance are higher than that of comparable conversion layers (DOW and HAE). Correspondingly sealed layers withstand approx. 600 hours in the neutral salt spray test according to DIN 50 021 SS. On an initial thin layer, which is in direct contact with the metal substrate, there is a slightly porous oxide ceramic layer. This layer carries a very porous oxide ceramic layer of equal thickness. The second layer may serve as an adherent base for paint or impregnations such as, for example, PTFE or other lubricants, which leads to improved corrosion and anti-friction properties. All commonly used magnesium alloys can be treated with the MAGOXID-COAT® process. Tests carried out with AZ91HP showed that the MAGOXID-COAT® treatment has only an insignificant negative effect on the fatigue strength of the base material. The MAGOXID-COAT® surface treatment is suitable for many industrial applications where magnesium substrates are used under special conditions as, for example, in mechanical engineering and the automotive sector. From the many varied applications we give here two as examples. ®

532 One possibility for the marking and signing of packaging materials is the application of hot stamping. A manufacturer of pharmaceutical products uses printing plates made of magnesium. Usually, brass or steel printing plates are used for the hot stamping of packaging materials. The printing letters are milled, engraved or ground and may be changed according to application. The manufacture of such printing clichés is comparatively expensive and is only profitable for a large quantity of, for example, one million impressions. For smaller quantities, plating plates made from a magnesium die-casting alloy have become of interest now, typically with a 100 x 50 mm size. The printing letters are manufactured by etching from defined phototechnically treated areas of the magnesium Picture 4: Magnesium printing plates, plate. This method of production is more economic than treated with 20 μm MAGOXIDthe conventional mechanical processing of brass or COAT steel. On the other hand, however, magnesium-based (Photo: AHC) clichés only have a short lifetime if not surface treated. Particularly when used for applications in the pharmaceutical sector the cleaning of the magnesium surface with disinfectants leads to a certain roughness that in turn has a negative influence on the printing image. A surface treatment ® with a 20 μm MAGOXID-COAT layer (picture 4) offers a solution to the problem. It is also possible to produce black layers that contain fade-resistant and chemically inert spinels. It is possible to treat almost any magnesium substrate suitable for industrial use. ® One intended application of MAGOXID-COAT black is, for example, the inner coating of optical components and precision threads, but is also suited for application to heat radiators, in the vacuum technology and the aviation and aerospace industry. A well-known sensor manufacturer has developed a light and ultra-compact optical sensor, destined especially for the automobile industry. This non-contact sensor measures the longitudinal and transversal vehicle dynamics relative to the road surface, thus providing data such as transverse angle, longitudinal and transverse speed as well as longitudinal and transverse acceleration. Due to its compact dimensions (approx. 164 x 52 x 61 mm) and light weight of 500 g this sensor is especially well suited for mounting on the vehicle wheel (also when being steered) for tyre-slip angle measurement. In order to be able to realise this compact and lightweight design a carbon fibre reinforced plastic has been used for the outer housing of the sensor whilst the engineers have decided to use a magnesium alloy for the housing containing the optics, as well as for further components (picture 5). These components require an adequate corrosion protection together with a black surface that reduces light reflection within the optical path. ® These requirements are met by a deep black, 25 μm MAGOXID-COAT layer. Correspondingly sealed layers withstand approx. 400 hours in the neutral salt spray test according to DIN 50 021 SS. Due to its electric insulation properties it effectively avoids ® contact corrosion. The degree of reflection of a black MAGOXID-COAT layer is less than 5%. ®

533 Today this optical sensor with the black magnesium components has proved to be excellent in practice and contributes considerably to precise, multi-dimensional vehicle measurement engineering. A surface coating well suited to treat magnesium alloys will certainly contribute to an increased application of this most interesting substrate used for structural purposes. However, even the best surface protection is of poor efficiency if the magPicture 5: Optical vehicle sensor: The CFRP outer nesium substrate does not fulfil certain housing contains magnesium components coated with basic requirements. Such as, for example, a deep black MAGOXID-COAT layer from AHC a high-purity quality of the base material, a Oberflächentechnik. magnesium cast that is almost free of (Photo: CORRSYS-DATRON GmbH) pores, or the avoidance of surface impurities with blasting material. A successful surface treatment requires an extensive dialogue between the moulder, the coating company and the user. ®

Literature [1] Kurze, P.: Chromfreie passivierte Magnesiumoberflächen mit nachfolgendem Lackauftrag. 8. Magnesium Abnehmerseminar / 8. Magnesium Automotive Seminar der Europäischen Forschungsgemeinschaft Magnesiumguss e.V. (EFM), 14./15. Juni 2000 in Aalen [2] EP 0 333 048 [3] Olbertz, B.; Haug A.T.: Oberflächenschutz für Magnesiumwerkstoffe, Metalloberfläche 43 (1989) 4, S. 174–178 [4] Kurze P.: Electrochemical Coating of Magnesium Alloys, Tagungsmaterial der DGM Magnesium Alloys and Their Applications, Garmisch-Partenkirchen, April 1992 [5] Kurze P.: Magnesiumlegierungen elektrochemisch beschichten, Metalloberfläche 48 (1994) 2, S. 104–105

534

Approach to Control the Corrosion of Magnesium by Alloying V. Kaese12, P.-T. Tai2, Fr.-W. Bach2, H. Haferkamp3, F. Witte4, H. Windhagen4 1

Volkswagen Group Research, Wolfsburg, formerly Institute for Materials Science and Biomedical Engineering Centre, University of Hanover 2 Institute for Materials Science, University of Hanover 3 Institute for Materials Science, Biomedical Engineering Centre, University of Hanover 4 Hanover Medical School, Hanover

Abstract For Magnesium alloys there are more fields as transport, communications or batteries. One is the use as bioresorbable implant in bone surgery. This saline area has high demands concerning corrosion resistance and is therefore used as main target for alloy development with the leading aspect of corrosion but regarding mechanical properties. A step-by-step building of alloy systems leads to an enhancement of the corrosion resistance by factors. These alloys are composed using common and newly defined principles of corrosion resisting alloying and processing parameters as extruding and heat treatment which aim for the enhancement of corrosion resistance. The final step is the validation of corrosive and mechanical characteristics of Fluorine-alloyed LAE-alloys in comparison to standard alloys as well in laboratory testing as in-vivo-implant in genuine pig [1].

Introduction and Motivation Table 1. Characteristics of Magnesium

Advantage availability density damping neutron holding cross section machining biocompatibility

In between corrosion casting characteristics

Disadvantage ductility creeping electrochemical position

Magnesium pros and cons are dominated by aspects as density and castibility vs. strength, creep and corrosion resistance Table 1. But the characteristics of Magnesium are more complex. The electrochemical border position can be controlled in the field of transportation if the junction is separated with coated bolts or washers [2]. Even the low corrosion resistance in saline environment is not only disadvantageous, it is wanted in the field of sacrificial anodes or batteries. Magnesium batteries are used not only to operate life vest lights or sonar buoys.

535 Experiments as implant material are dating back to 1906 [3]. According Wolf’s Law of Transformation best bone-building is achieved if the fracture is loaded. With increased bone-building the implant has to support less until relief at recovery. Best results will be gained if the implant has comparable mechanical properties as the recovering femur. The motivation to develop implants with sufficient mechanical properties is strong because there is no alloy on stock fitting with the mechanical properties of human femur Figure 1.

Figure 1. Mechanical properties of common implant materials, AZ31 vs. human bone

Besides the mechanical fitting, Magnesium is essential for some bodily functions as the nervous system. Over doses will be spilled out by the urinary tract. These boundary conditions allow further thinking about Magnesium as implant. Fractured collar bones are fixed with Magnesium splints, which lead to a sufficient recovery of the fracture and total dissolving of the implant however coming along with the formation of large gas volumes in vivo [3]. Magnesium dissolves in saline environment as salt spray in the streets or in saline solution of human body. The deductions are that future resorbable Magnesium implants, which could spare the often necessary operational revision, have to be less gas producing in the course of corroding, see [4], and more ductile and stronger. MgY1.6Cd0.25Ag0.3MnCa and MgNd2.46Cd0.12AlMnSi corrode with a hydrogen for3 2 3 2 mation of 0.8 cm H2/(cm ⋅ d) and 1.05 cm H2/(cm ⋅ d). These alloys produce far less 3 hydrogen as the theoretical maximum of 2.25 cm H2/(cm2 ⋅ d). The bone building and recovery is accelerated by factor 1.5 [5]. But the Cd-content is leading to dead-end seen from the aspects of machining and bioresorbable material. Some experiments have been performed as anode vs. a Pt-cathode in thorax of a dog. The corrosion rates of MgMn1.5, MgAl3Zn1 and MgAl6Zn3 vs. the Pt-cathode have been five times higher than necessary for a life of five years [6], see [1]. Besides orthopaedic surgery there are patented developments for heard-thorax solutions as expandable stent [7].

536

Theoretical Approach and Methods Discussed AZ31 or AZ63 corrode to fast and release to much hydrogen. The corrosion rate for in vivo corroding metal has to be controlled by alloying and metallurgical processing, not by coating, which could cause stroke. There is data to evaluate the influence of alloying elements but only partly sorted by principles. The metallurgical processing can be focused e.g. on controlling the impurities or to homogenise precipitation in the matrix. The influence of the discussed elements. Figure 2 can be summarised in principles for alloying components Table 2, see [1]. H Li

Be

B

C

Na Mg

Al

Si

Ca Sr

Cr Y

Zr

Mn

Fe

Co

Mo

Ni

Cu

Zn

Ag

Cd

Pr

Nd

Gd

Dy

O

F

S

Cl

Ga

La

Ce

N

Er

Sn

Sb

Pb

Bi

Yb

Figure 2. Periodic system of elements whose influence on Magnesium-corrosion is described

Some principles are written down now Table 3, as example #14 is chosen to be discussed here. Magnesium does corrode in NaCl-solution because the Mg(OH)2-layer is not stable to NaCl, the forming MgCl2 is soluble and not protecting. Looking for some alloying component which could form insoluble layer components, we choose Fluorine. It is commonly known as component of conversion not alloying, however Fluorine is not easy to alloy and MgF2 is not thermodynamically stable to NaF. The first obstacle can be bypassed by using a Fluorine salt with a second component known to enhance the corrosion resistance: AlF3. The alloying of AlF3 in AZ31 is leading to a reduction of the corrosion rate by factor 2 in artificial seawater. The second obstacle, looking for thermodynamically stable vs. NaF and water insoluble layers has to be gotten around by calculation. The enthalpy of reaction of NaF and Fluorine salts composed with known alloying components is compared, e.g. leading to LiF and CaF, both nearly insoluble in water solution Figure 3. Lithium is chosen, because it does fit better with the principles than Ca, which even causes hot cracking. Fluorine is successfully inserted in the hdp-LAE system by alloying [9].

537 Table 2. State of the art principles of corrosion influencing alloying of Magnesium, see [1]

1.

Avoiding elements with noble electrochemical potential or lower hydrogen overvoltage Avoiding water solvable inclusions as MgCl2, Mg3N2 Precipitating impurities as Al-Fe-compounds by Mn Grain refining of Mg-Al-alloys by adding carbon to reduce porosity Rapid solidification to enrich the matrix which corrosion resistance enhancing elements Dehydrating and compacting the natural MgO-Mg(OH)2-layer, e.g. by Al or Ce Binding the anions by opposite charge Building of amorphous, compacted and relaxed layers, e.g. by Cr, even to shift potential Alkalising the layer to shift pH-value over 9.5 by Li to stabilise the natural layer

2. 3. 4. 5. 6. 7. 8. 9.

Table 3. Now defined principles to enhance the Magnesium corrosion resistance by alloying

10. 11. 12. 13. 14. 15. 16.

Maximised solution of corrosion enhancing component in the matrix Geometrical fitting of the lattice of layer components Binding Mg-cations by double charged oxidised groups Thermodynamic calculation of multiphase micro structure with aspect of corrosion Thermodynamic calculation of layer stable vs. NaCl-containing solutions Absorption of neutrons (for using Magnesium in atomic power plants) Biocompatibility of components

Enthalpie of reaction [kJ/mol]

-300 SnF2 PbF2

-500

ZnF2

-700 2/3 AlF3

MgF2

-900

2 LiF

-1100

CaF2

-1300 0

100

200

300

400

500

Temperature [°C]

Figure 3. Enthalpy of reaction of NaF and Fluorine salts, data see [8]

600

700

800

538 Results Gastaschen hydrogen bubbles

Detail LAE442

Detail

LAE442(AlF3)2 at%, ϕ = 3,7

Figure 4. Extruded LAE442 (top) and LAE442(AlF3)2 at, X-ray after 2 weeks of genuine pig femur, width of picture: 120 mm

The effectiveness of Fluorine in LAE442 as cast is reached with 2at% added to the Magnesium-content. The corrosion rate of as cast material in potentiostatic dipping test in artificial seawater acc. to ASTM B 1141–51at room temperature is coming down from 0.2 mm/a to 0.15 mm/a to 0.1 mm/a to 0.1 mm for 0, 1, 2 and 4 at% (AlF3). LAE442 plus 2at% AlF3 is tested in extruded condition to enhance mechanical properties in comparison with LAE442 in vivo in genuine pig femur and shows for the reference LAE442 bubbles of hydrogen and signs of disintegration. LAE442 plus 2 at% (AlF3) shows original geometry and no visible hydrogen bubbles Figure 4.

Conclusions and Outlook 16 principles of alloying Magnesium to enhance corrosion resistance are proposed, one example is presented and jointed with the task to design an alloy with high corrosion resistance for Magnesium implants, one of the oldest fields of Magnesium applications. It is shown that with thermodynamic calculation the corrosion resistance by Fluorine addition can by enhanced further in synthetic seawater and in vivo in genuine pig. The reference LAE442 is to expensive for common use, but it has perspective in the field of implants but alloying with Fluorine containing salts will reach other alloy systems.

539

Acknowledgement The research has been supported by the German Ministry of Education and Research BMBF as one winning project named RemOs of the Program KOMED in 1999 and the Deutsche Forschungsgemeinschaft, Program SFB390. Special thanks to P. Lindner and P. Juchmann, Salzgitter, to A. Pisch, Grenoble, and R. Schmid-Fetzer, Clausthal, for thermodynamical advise plus calculating and the work group of RemOs.

References [1] V. Kaese, Beitrag zum korrosionsschützenden Legieren von Magnesiumwerkstoffen, VDI-Verlag, Düsseldorf, 2002. [2] H. Schreckenberger, Korrosion und Korrosionsschutz von Magnesiumwerkstoffen für den Automobilbau – Problematik der Kontaktkorrosion, VDI-Verlag, Düsseldorf, 2001. [3] A. Lambotte, Bulletins et Mémoires de la Société Nationale de Chirurgie Vol. 58 (1932)2, p. 1325-1334. [4] M.G. Seelig, Archives of Surgery Vol. 8 (1924) p. 669-680. [5] G.B. Stroganov et al., German Patent DE 1953241, 1971. [6] G. Fontenier, R. Freschard, M. Mourot, Medical and Biological Engineering Vol. 13 (1975)5, p. 683-689. [7] H. Haferkamp, V. Kaese, M. Niemeyer, M. Rodman, Tai Phan-tan, B. Heublein, R. Rohde, Magnesium 2000, Editors E. Aghion, D. Eliezer, Dead Sea, Israel, p. 159-164. th [8] O. Kubaschewski, C.B. Alcock, P.J. Spencer, Materials Thermochemistry. 6 ed. Oxford, Pergamon Press, 1993. [9] P. Juchmann, Beitrag zur technologischen Eigenschaftserweiterung von MagnesiumWerkstoffen durch Lithium, VDI-Verlag Düsseldorf, 1999.

540

Magnesium Alloys Protection by Microplasmic Processing Belozerov Valeriy, Isakov Sergey, Makhatilova Anna, Subbotina Valeriya National Technical University “Kharkov Polytechnic Institute”, Kharkov; Isakova Tatiana – Innovation Center “MAGIC Solutions” ltd., Kharkov

Technology description Magnesium alloys have high reactivity and at maintenance in atmospheric conditions can corrode with a noticeable velocity. The base modes of Magnesium alloys protection from corrosion are both oxide layers formation and protective coatings deposition. However, the application of traditional modes of coating deposition are not always reduced to required properties reaching. Last years new methods of coating deposition and surface hardening with use of high concentrated power sources intensively explicate: electric discharges, lasers etc. One of such methods is the microplasmic processing (microarc oxidation), which permit to get functional coatings on a detail surface. The both deformable AM60B (Mg - 6% Al - 0,4% Mn) and cast AZ91 (Mg - 9% Al 0,7% Zn - 0,3% Mn) Magnesium alloys were processing by MPP. The process was carried out in anodic-cathodic condition in an alkaline electrolyte with adding of inorganic compounds. The power supply of a condenser type was utilized. It is necessary to mark, that it was impossible to organize microarc discharges on Magnesium alloys surface in an alkaline electrolyte. It is bound that on a surface of an alloy the dielectric barrier layer is not formed because of lack of interaction of Magnesium with alkali. Use of water solution of a silicate (NaSiO3) has allowed to transfer process in arc discharges conditions. The power of such discharges is very high, that there are both melting of metal and formation of craters. The steady microplasmic process was implemented in composite electrolytes containing the following compounds NaOH, KOH, NaAlO2, Na5P3O10, NaSiO3 etc. As a result of such processing the surface layers of Magnesium alloys will be converted to a ceramic coating consisting of crystalline oxides and spinel. The X-ray phase analysis has shown, that the base phases of a coating are MgO, MgAl2O4, MgSiO3, the quantitative relation between which is defined by both electrolyte structure and electrolysis parameters (processing duration, current density and other). Thus, the organization of process in microarc discharges conditions requires the relevant selection of an electrolyte structure and electrolysis parameters, that allows to form coatings by thickness 200-300 microns at overgrowth velocity 1-2 microns per minute. The coating hardness makes HV 500-600 that is at 8-10 time higher then the hardness of a base (HV 60).

541 Two structures of an electrolyte distinguished by a specific resistance and acidity (pH) were utilized and also the following parameters were varied: – Current density from 5 up to 20 A/dm2; – Duration of processing from 4 about 120 minutes; – Electrolyte structure.

Test procedures and results Surface roughness There are no difference between surface roughness of AM60B and that of AZ91 alloys. Surface roughness increases with MPP coat thickening. Surface roughness varied from 9 to 20 microns (Rz) with coating thickness from 9 to 40 microns. Rz means 10-point average peak height difference of surface roughness. Surface roughness measured by roughness measuring equipment. SEM observation of the samples shows the good surface roughness. Adhesive test The microplasmic processing ensures a high adhesion of a coating with a base. At the first stage the treated surface was scratched cross-hatching 10 x 10 lines of 1mm interval on the surface by a sharp knife and then were peeled by adhesive tape and checked whether at least one cross-hatching area (1x1mm) is peeled or not. All tested coatings were not peeled from substrates. At the second stage cross-hatching samples were dipped into water for 240 hr and then were peel by adhesive tape and checked whether at least one cross-hatching area (1 x 1mm) is peeled or not. All tested coatings were not peeled from substrates. Complex corrosion test X-figurative cross-hatching samples were immersed in complex corrosion cycle bath for a certain long time (2 months). The corrosion cycle includes salt spraying (1 hour) and drying step (1 hour) and so on. After that the corrosion width were measured at the crosshatching lines. The following conditions were used: – the salt solution: 5mass% NaCl (95mass% water) – temperature of the exposed zone: 50 degrees in C – the pH range from 6.5 to 7.2 The criteria are as follows: Fair/ F: Corrosion width x ≤ 0.5 mm Marginal/ M: Corrosion width 0.5 mm < x ≤ 1 mm Poor/ P: Corrosion width 1 mm < x ≤ 5 mm Very Poor/ VP: Corrosion width x > 5 mm The results have shown Table 1, that in case of a cast alloy AZ91 (Mg - 9% Al - 0,7% Zn - 0,3% Mn) at particular electrolysis conditions the corrosion sign was not revealed (corrosion width x ≤ 0,5 mm), whereas under other requirements of an electrolysis the corrosion width reached 1 - 5 mm.

542 On a deformable alloy AM60B (Mg - 6% Al - 0,4% Mn) was impossible to ensure a high corrosion stability of a coating Table 2. Table 1. Alloy AZ91

№ of a sample

AZ2 AZ4 AZ5 AZ6 AZ7 AZ8 AZ9 AZ12 AZ13 AZ14 AZ21 AZ15 AZ18 AZ22 AZ16 AZ24

Mode of MPO Type of an electrolyte

Current density, A/decimeter2

Time, mines

1 1 1 1 1 1 1 1 1 2 2 2 2 2 2 2

10 20 20 10 10 20 20 40 5 20 20 20 20 20 20 20

30 20 20 15 15 10 10 9 120 15 15 8 8 8 4 4

Thickness, micron

Criteria

23-27 29-32 28-33 14-15 14-16 13-16 15-17 22-26 39-44 24-26 20-24 16-19 13-17 8-12 12-14 15-17

M M P M F F M M M P M P M M M M

Thickness, micron

Criteria

23-27 31-33 32-36 14-16 17-20 18-20 23-27 35-36 21-24 16-18 8-10 13-16

P P P P VP P P P P P VP P

Table 2. Alloy AM60B

№ of a sam- Mode of MPO ple Current density, Type of an A/decimeter2 electrolyte

Time, mines

AM2 AM4 AM5 AM6 AM8 AM11 AM13 AM15 AM30 AM28 AM18 AM19

30 20 20 15 10 10 9 90 15 8 4 4

1 1 1 1 1 1 1 1 2 2 2 2

10 20 20 10 20 20 40 5 20 20 20 20

543 The corrosion trials have shown, that the protective properties depend on coating phase structure, which defines its specific volume. In that case, when the ratio of coating specific volume to the base one is more then 1, the formed coating is under compression stresses operation, that reduces in improving its continuousness and increase of a corrosion stability. If the ratio is less then1, the coating is under strain stresses operation, that reduces its corrosion stability. To get high anticorrosive properties of a coating it is necessary to increase a content of spinel MgAl2O4 that is achieved by Aluminum content magnification in an alloy or in an electrolyte. Thus, MPP technology can successfully be used for a raise of a corrosion stability and hardness of magnesium alloys.

544

Class A Surface Quality for Magnesium Diecastings Using Ceramic Precursor-based Coatings R. Gadow, D. Scherer, C. C. Stahr University of Stuttgart, Institute for Manufacturing Technologies of Ceramic Components and Composites – IFKB, Stuttgart, Germany

1

Introduction

Weight reduction is today one of the priority objectives for the automotive industry. More and more safety and comfort systems are integrated to modern vehicles. This means an increase of weight and therefore an increase of fuel consumption. In order compensate this and to reduce environmental pollution automotive manufactures are forced to implement light weight design in every component of a new car. Magnesium (Mg) provides a further weight reduction compared to other light metal alloys, concerning to its low specific weight which is still 30% under Aluminum. A major part of magnesium is manufactured by high pressure diecasting process. Major drawbacks of Mg are the poor surface properties concerning corrosion and friction as well as the fire risk during machining operations. To make Mg-parts useable for the automobile body a well designed coating is needed which covers corrosion protection and optical requirements in “class A quality”.

2

Painting systems for automotive bodies

Paintings for automotive bodies need not only to protect the metallic substrate before corrosion. They have additionally to match high decorative requirements. To fulfill these requirements a multilayer system is needed. Although many studies have been made to prevent Mg alloys from corrosion, no system is well established yet especially not for the automobile body. The more promising methods are hindered by fairly high production costs. Normally the first step to improve corrosion resistance of a light metal is a conversion coating. This improves paint adhesion and prevents the spread of corrosion under the paint film [1]. The commonly used chromium containing baths are no longer of interest for new coating systems due to their critical environmental and health hazard. Two environmental friendly chromium free conversion coatings are commercially available, a fluorozirconate and a phosphate-permanganate treatment. Details of this conversation coatings are given elsewhere [2], [3], [4]. Second step is a primer which has the function to protect against spread of corrosion and mechanical damage as well as to reduce the surface roughness [5]. In the third step a filler is applied. The filler improves mechanical resistance against chipping as well as it increases the resistance against moisture. The filler may be grinded or polished to improve optical appearance. At last the top coat is applied which can be a one layer solid color coat or a metallic two layer coat.

545

3

Painting systems for magnesium parts

On the one hand there are the optical requirements like excellent gloss and smoothness for visible outer parts in automotive applications, on the other hand Mg diecastings have very poor surface properties. Mg can only be very costly grinded due to its reactivity. Therefore a new coating is needed which can smooth the surface without mechanical posttreatment. Of course the paint system must protect reliably against corrosion. To minimize investment costs the new system needs to be integrated into available standard coating lines. The first step is again a chromium-free conversion coating, which is explained in chap. 2. As second layer a silicon-modified organic resin, so called precursor, is applied. These resins excel advantageous curing conditions, good adhesion and good mechanical properties. The silicon part improves hydrophobicity and weather resistance, UV stability and temperature resistance. Also the silicon-modified resins have excellent spreading and leveling behavior [6]. Last but not least they can be applied through various standard processes like dip-coating, air-spraying, electrostatic spraying and brushing. In that they can be manufactured in available coating lines. To finish the coating system and achieve the final optical appearance a standard one or two layer topcoat is applied. Fig. 1 and Fig. 2 show schematically a standard coating on Mg substrate compared with a precursorbased coating. Clearcoat (20-30 µm)

Clearcoat (20-30 µm)

Basecoat (10 µm)

Basecoat (10 µm)

Filler (50-120 µm) Wet paint or EPS

Precursor (30-70µm)

Primer

Conversion coating

Chromate Magnesium

source: Norsk Hydro

Figure 1. typical coating used for magnesium wheels

4

AM 50 substrate

Figure 2. coating system for magnesium diecastings in Class A quality

Corrosion Test

The corrosion behavior of the precursor-based coating system was tested by salt spray according to ASTM B117 (240 h) and cyclic testing according to VDA 621-415 (10 cycles). Three different AM50A samples were pretreated with different conversion coatings: chromate, fluorzirconate and phosphate-permanganate. The following precursor layer was applied by pneumatic air spraying and cured at 160 °C for 20 min. To complete the coating system a commercially available two layer top coat was applied: a metallic basecoat and a clearcoat. The samples were scribed to the metal before the corrosion testing. The results of salt spray and cyclic testing are shown in figure 3 and 4.

546

phosphate permanganate pretreated

fluorozirconate pretreated

chromate pretreated

Figure 3. Coating systems after 240 h ASTM B117 salt spray testing with scribe

cyclic test VDA 621-415

phosphate permanganate pretreated

cyclic test VDA 621-415

fluorozirconate pretreated

cyclic test VDA 621-415

chromate pretreated

Figure 4. Coating systems after 10 cycles according to VDA 621-415

No matter which conversion coating was used, all coating systems showed excellent corrosion resistance in both tests. None or only minimal scribe creepage can be detected [7].

547

5

Conclusion

Precursor-based coating systems open the way for magnesium in the automobile body. Due to the possibility to integrate the new coating in available standard coating lines it can economically be deposited. Posttreatment of the diecastings is minimized and a class A surface finish is achieved. Sufficient corrosion resistance can be achieved even if a nonchromate pretreatment is used.

6

References

[1] K. Yasuhara, “Surface pretreatment of aluminum for automotive applications” in Automotive Paints and Coatings (Ed.: G. Fettis), VCH Verlagsgesellschaft, Weinheim, 1995, pp. 9-27 [2] J. I. Skar, M. Walter, D. Albright, “Non-chromate conversion coating for magnesium die castings” SAE technical paper series 970324, 1997 [3] H. Gehmecker, “Chrome free pre-treatment of magnesium parts” in Proceedings of IMA 51, 1994 May 17-19, Berlin, Germany, pp. 32-34 [4] D. Hawke, D.L.Albright, “A phosphate-permanganate conversion coating for magnesium” in Metal Finishing, Elsevier Science, 1995, pp. 34-38 [5] Z. Vachlas, “Primers for the automotive industry”, in Automotive paints and coatings (ED.: G. Fettis), VCH Verlagsgesellschaft, Weinheim, 1995, pp. 28-71, [6] R. Gadow, F.J. Gammel, F. Lehnert, D. Scherer, German Patent application DE 10006270.9 [7] R. Gadow, F.J. Gammel, F. Lehnert, D. Scherer, J. I. Skar, „Coating systems for magnesium diecastings in Class A surface quality“ in Magnesium Alloys and their Applications (Ed.: K. U. Kainer), WILEY-VCH Verlag GmbH, Weinheim, 2000, pp. 492-498

548

Wear and Corrosion Resistant Ceramic-based Multilayer Coatings with dry Lubrication Ability R. Gadow, D. Scherer, C. C. Stahr University of Stuttgart, Institute for Manufacturing Technologies of Ceramic Components and Composites – IFKB, Stuttgart, Germany

1

Introduction

Magnesium alloys offer new ways in light weight engineering due to their low specific weight and their high strength to weight ratio. The use of magnesium alloys in many engineering applications is mainly limited by their unsatisfying surface properties and the poor corrosion resistance. Therefore an effective and appropriate coating technology has to be adapted to enable a successful implementation of magnesium alloys under tribological load and in corrosive environment. In most cases a one layer coating is not able to meet all demands. To match the different requirements multilayer coatings have to be used. The coating is combined of at least one hard coating which protects against wear and an additional coating ensuring functionality and improved dry lubrication ability. This can either be a polymer or a thin film.

2

Coating technology and experimental setup

Several approaches to increase the wear and corrosion resistance were done before. Lacquers for example may protect against corrosive attack, but only as long as they are not hurt. They also may provide dry lubrication if filled with dry lubricants like PTFE, but abrasion resistance is normally low. Thin film technology offers coatings with e.g. dry lubrication ability, which are hard and abrasion resistant. But normally they are brittle and due to the low Young’s modulus of light metals they crack if load is increased too much. Thermally sprayed hard coatings resist reliably against wear, but due to process-related coating porosity corrosive environment may attack through the coating. Thermally sprayed coatings may provide dry lubrication ability in some cases, but normally the coefficient of friction is high. The solution must be combined multilayer coatings, consisting of a hard thermally sprayed layer and an additional lacquer or thin film layer [1]. For the hard coating ceramics, metals and cermets can be used. The decision what kind of material is to be used depends mainly on the application. To show the tribological behavior of combined coatings a TiO2 primary coating was applied to a magnesium substrate by Atmospheric Plasma Spraying (APS). For details of thermal spray technologies see [2]. It was covered on the one hand with an a-C:H thin film. On the other hand a dry lubricating varnish containing PTFE and MoS2 was applied by air spray. The lacquer was applied to the TiO2-surface once as sprayed and second on a grinded surface. If a thermally sprayed

549 coating is combined with a polymer coating the process-related roughness is useful to keep the dry lubricant filled lacquer on the surface. Even if the peaks are worn dry lubricants remain on the ground. Thin films are only applied to a grinded surface. The hard coating supports against cracking under high load, peaks would be a potential source of failure. The tribological results were made by a pin-on-disc test with a normal load of 10 N, constant humidity of r.h. = 40% and temperature of 25 °C. The counterpart was a hardened steel 100Cr6 ball with a diameter of 5 mm. In a second investigation basic corrosion tests were done with thermally coated Mgsubstrates. In order to evaluate the behavior of different materials a metal (Al/Si), a cermet (WC-CoCr) and two oxide-ceramics were applied to AZ91 Mg-plates. The coatings were done by APS-spraying except for the WC-CoCr which was done by High Velocity Oxygen Fuel (HVOF) Spraying. To expose only the coated side of the Mg-plate a tube was glued to the surface. The tube was filled with solution of NaCl (3%). After 24 h the tube was removed and the surfaces was optically examined.

3

Results and discussion

3.1

Tribological results

The hardness of the TiO2 coating was measured in the cross section and was HV 0,5 = 962. Roughness as sprayed was Ra = 3.7 μm. The grinded surface roughness was Ra = 0.05 μm. Figure 1 shows the coefficient of friction versus the number of oscillations for the different coating systems. As a reference the data from a TiO2 coating without functional second coating is added. 1

TiO2 not coated

Coefficient of friction µ

0.8 lubricant lacquer grinded

0.6

0.4

0.2

DLC

lubricant lacquer as sprayed

0 0

20000

40000

60000

Number of oscillations Figure 1. Tribological behavior of coated Mg substrates

80000

1000

550 Figure 2 shows the wear rate for the different friction pairings. The wear of the coating was measured through tactile surface scanning and a software-based loss volume calculation. The wear on the ball was calculated after measuring the spherical loss. 0.3 surface wear volumetric wear [mm³]

0.25

wear on 100Cr6 ball

0.2 0.15 0.1 0.05 0 DLC

lubricant lacquer; lubricant lacquer; TiO2 not coated TiO2 polished TiO2 as sprayed

dry lubricant coating Figure 2. Volumetric wear of substrate and counterpart in pin-on-disc test

The coefficient of friction for the uncoated TiO2 is not acceptable high, but anyway it is tolerably constant. Therefore no galling effects are expected. Wear on the coating surface is high and on the ball it is even the highest. For the combination grinded TiO2/lubricant lacquer a significant improvement of the friction coefficient is visible, but only in the beginning. After a short period the lacquer is worn and again TiO2 slides against 100Cr6. This results in high wear on the coating surface, but wear on the ball is reduced. The two systems grinded TiO2/thin film and TiO2 as sprayed/lubricant lacquer show excellent behavior. The coefficient of friction is constantly low. Wear is reduced to a minimum [3]. 3.2

Corrosion test results

Figure 3 to 6 show the results of the corrosion test after the removal of the tube.

551

Figure 3. Al2O3/TiO2-coating after corrosion test

Figure 4. TiO2-coating after corrosion test

Figure 5. Al/Si-coating after corrosion test

Figure 6. WC-CoCr-coating after corrosion test

The most severe corrosive attack is seen for the Al2O3/TiO2 and the TiO2 coatings. Both coatings show severe cracks and are delaminated from the Mg substrate. This is due to the high porosity of these APS sprayed coatings, especially in form of open porosity. Hydrogen is formed and pressure is build up to a certain degree [4]. As a result delamination of the ceramic coatings occurs. Large areas of the Al2O3/TiO2 and the TiO2 coatings delaminate simultaneously. For the AlSi coating instead only minor corrosive attack is seen at the end of the test changing the optical appearance of the area of the Al/Si coating which was exposed to the NaCl solution. This is due to the lower porosity of this coating and the quasi ductility of the Al/Si coating preventing a brittle failure, as in the case of the ceramic coatings. For the HVOF sprayed WC/Co/Cr coating severe local corrosive attack can be observed at areas of open porosity due to the high differences in the electrochemical potential. But due to the higher adhesive strength and the higher ductility of the HVOF sprayed cermet coating as well as due to the lower porosity a delamination of larger areas of the coating is prevented. Combined coatings were not tested at this time, but it is expected that corrosive attack will not occur as long as the polymer film is intact and a total electrochemical isolation is achieved.

552

Conclusion Combined coatings based on ceramic hard coatings are a powerful instrument to protect light metal and in this study particularly magnesium against wear. Functional polymer or thin film coatings enable dry lubrication ability. Ceramic hard coatings do not always protect against corrosion itself. Additional treatments and coatings are needed. The right combination of hard and functional coatings enables the application of light metal components in many fields. Automotive, aerospace or consumer products are only three of them. The costs of a combined coating are tolerably low particularly for the combination of a hard ceramic coating with a lubricant lacquer. The combination with thin films may be economical for small high end components.

References [1] R. Gadow, A. Killinger, D. Scherer, Ceramic Polymer Composite Coatings, Mat.wiss. u. Werkstofftechnik, 29, 1998, pp. 292-299 [2] L. Pawlowski, The science and engineering of thermal spray coatings, John Wiley & Sons Ltd. Chichester, 1995 [3] R. Gadow, D. Scherer, Novel processing for combined coatings with dry lubrication ability, Ceramic Transactions, Vol. 129, Innovative Processing and Synthesis of Ceramics, Glasses and Composites V, eds. N.P. Bansal et. al., The American Ceramic Society, Westerville, Ohio, 2002, pp. 125-136 [4] C. Friedrich, R. Gadow, D. Scherer, Functional Ceramic and Metallurgical Coatings on Magnesium Components, Proceedings of the International Thermal Spray Conference & Exhibition, 28-30 May 2001, Singapore

553

New Generation of Keronite Process for Corrosion Protection of Magnesium S. Hutchins, P. Shashkov, V. Samsonov and A. Shatrov Keronite Ltd, Granta Park, Great Abington, Cambridge CB1 6GP, UK

1

Introduction

Magnesium’s favourable strength-to-weight ratio makes it attractive to automotive and other industries. The cost of both raw magnesium and finished-goods is reducing as production increases. However, the lack of an acceptable-cost, chrome-free corrosion pretreatment has hampered its widespread adoption. This paper introduces a new generation [1] of plasma electrolytic oxidation (PEO) coating process that addresses the combined requirements of corrosion protection, scratch resistance, paintability (including A-class surfaces), acceptable cost, series production and large parts (up to 2 m2) capability. A further benefit of smoother coatings is to enable wear applications without the need to polish. This new PEO application category is mentioned here for completeness only and is not the main subject of this paper.

2

Process Description

The Keronite process is an environmentally-friendly, bath-based treatment for light alloys where a pulsed voltage is applied to the substrate component and the resulting sparks create a fused ceramic layer on the surface of the metal. The latest advancements in the Keronite process focus on optimisation of waveform and frequency of electrical pulses together with the introduction of acoustic vibration of a certain frequency range and air micro-bubbles into the electrolyte. Amplitude, duration and shape of both positive and negative pulses were optimised to achieve controllable high deposition rates of 2-10 μm/min. Acoustic vibration was introduced into the bath by pumping electrolyte through a specially designed hydrodynamic device working in the audible frequency range. Overlapping of electrical and acoustic frequency ranges had a synergistic effect on the process. The hydrodynamic generator is also responsible for oxygenating the electrolyte. The resulting Keronite process produces hard, smooth and compact layers on Magnesium.

3

Coating Characterisation

3.1

Morphology

Traditional PEO coatings are porous and have roughness increasing with thickness. This gives a good key for topcoats but limits hardness, corrosion performance of unsealed

554 specimens, and achievement of an A-class paint finish. Samples of AZ91D alloy were coated with 10 μm of Keronite and viewed by Electron Microscope. The surface has a lenticular structure with fine pores (Figure 1), smaller than with anodised Mg, also appears smoother and denser [2]. These characteristics will help roughness and hardness issues whilst allowing enough porosity for paint adhesion. X-ray analysis shows that the main constituent is spinel (MgAl2O3).

Figure 1. SEM of surface of 10 μm Keronite layer on AZ91D

3.2

Hardness

The in-plane hardness of a 60 μm Keronite layer was tested at Cambridge University [3] as a function of distance through the coating using a Berkovitch nanoindenter with 10 mN load (Figure 2). The average hardness of 7.1 GPa (approx 700 HV) compares favourably with that of bare magnesium (approx 100 HV) and is similar to that of hardened tool steel. 16

10 mN load, Berkovitch indenter

14

Hardness: 7.1±2.9 GPa

Hardness /GPa

12

10

8

6

4

2

Top of Coating

Substrate Interface 0

0

10

20

30

40

50

Approximate through thickness scale/μm

Figure 2. Hardness profile of 60 μm Keronite coating on AZ91D

60

70

555

4

Corrosion Properties

4.1

Description of Samples

Plates of cast AZ91D were coated with nominal 10 μm of Keronite. The authors found that 10 μm provides a balance of corrosion protection and scratch resistance with acceptable cost and coating time (2 minutes). The coating is designed as a pre-treatment for other sealers or top-coats. Therefore, the samples were post-treated with different coatings, including automotive-standard e-coat and powdercoat [4], as shown in Table 1. Automakers seek protection against galvanic corrosion and corrosion after the coating has been scratched. Sample plates were prepared accordingly. Table 1. Description of AZ91D samples prepared for testing.

Sample Label

Description

A.1 – A.4 B.1 – B.4 C.1 – C.4 D.1 – D.4 E.1 – E.4 F.1 – F.4 G.1 – G.4

10 μm Keronite as-coated 10 μm Keronite plus sodium silicate sealant (Na2SiO3) 10 μm Keronite plus sol-gel sealant (NTC GmbH, Basecoat U-Sil100) 10 μm Keronite plus e-coat1 (McDermid Electrolac High Build XD4434) 10 μm Keronite plus e-coat2 (BASF GV82/9438) 10 μm Keronite plus powder-coat (HB Fuller P4M5229 polyester) 10 μm Keronite plus e-coat2 plus powder coat

H.1 – H.2

Bare Magnesium

4.2

Corrosion and corrosion after scratch

One sample of each type (A.1, B.1, etc) was scribed through to the substrate and subjected to 750 h of salt spray according to ASTM B117. The corrosion was rated according to ASTM D1654 (Figure 3, 4 Table 2). All coatings show ratings of 9 or 10 after 750 h, in both scribed and unscribed regions. The presence of e-coat or powder coat (D.1-G.1) ensures a 10 rating in both categories. These results exceed those after 240 h of Beals et al. [5]. 10

9

8

Rating (ASTM D1654)

7

6

Keronite K + silicate sealant K+ sol-gel

5

K + e-coat1 K + e-coat2

4

K+ powder K + e-coat2 + powder

3

Mg

2

1

0 0h

24h

96h

192h

336h

500h

750h

Time in salt spray (h)

Figure 3. Corrosion resistance of Keronite plus additional coatings on AZ91D according to ASTM B117

556 10

9

8

7 ASTM 1654 Rating

Keronite K + silicate sealant

6

K+ sol-gel K + e-coat1

5

K + e-coat2 K+ powder 4

K + e-coat2 + powder Mg

3

2

1

0 0h

24h

96h

192h

336h

500h

750h

Time in Salt Spray

Figure 4. Corrosion creep from scratch using ASTM B117 test and ASTM D1654 corrosion creep rating.

4.3

Galvanic corrosion

One sample of each coating (A.2, B.2, etc) was drilled before coating and fitted with a zincplated M6 steel bolt after coating. The bolt was fixed using a nylon nut and the samples subjected to 750 h salt spray. The intention with these samples is to continue the test until 2000 h and rate using ASTM D1654 with the bolt removed. Nonetheless, inspection of the samples at 750 h was very encouraging on samples with e-coat or powder coat (D.2–G.2). 4.4

Phospahte Bath Compatibility

A wish of some automotive Mg users is to paint components in a conventional steel paint line. The early stages of such a process usually include zinc phosphate then e-coat. Magnesium is corroded by phosphating and cannot be painted on such a line. Samples A.3, B.3 etc were coated in full, dipped in zinc phosphate (ABPHOS 700Z) at Plastic Coatings Ltd, then tested by ASTM B117-97 with only the as-coated Keronite showing any sign of corrosion after 750 h (Figure 5, Table 2). Other workers [6] claim success passing Keronite-coated AZ91D through a steel paint process, including zinc phosphate, with no subsequent corrosion problems.

5

Mechanical Performance

5.1

Scratch Resistance

Samples A.4–G.4 were tested at the University of Hull [7] for scratch resistance using a Rockwell-type conical diamond indenter (tip Ø = 0.4 mm) in an IonCoat Multipath Scratch Tester to produce 4 parallel scratches with linearly increasing normal load Ln, up to 100 N. Table 2 shows the load (Lc2) that detached the composite coating from the substrate. As-coated Keronite detached at Lc2 = 14 ± 2N (approximately 3 times higher than anodised Mg [6]). The two sealants increased the resistance approximately 30% and e-coat1 by nearly 3 times. For coatings (F.4, G.4) including powder coat, the indenter did not detach the Keronite, even with maximum loading.

557 10

9

8

ASTM 1654 Rating

7

6

5 Keronite

4

K + silicate sealant K+ sol-gel

3

K + e-coat1 K + e-coat2

2

K+ powder

1

0 0h

24h

96h

192h

336h

500h

750h

Time in Salt Spray

Figure 5. ASTMB117 corrosion performance after immersion in zinc phosphate bath

5.2

Paint Adhesion

Samples A.4-G.4 were sprayed on one half with Newlon N17 water-based paint, and on the other with Trimite B401AE 262/3 acrylic solvent-based paint. These were then subjected to and rated by paint adhesion test ASTM D3359-97. The results (Table 2) show that the as-coated Keronite provides an excellent key for the paint. On samples B.4-G.4, any delamination occurred between the top two layers of coating, leaving the Keronite/polymer interface intact. Table 2. Summary of corrosion and mechanical performance of Keronite on AZ91D

750 hour Salt Spray Corrosion Resultsb

A B C D E F G H

Scratch Resistance

Paint Adhesion Ratingf

Coatinga

Subjective Creep After Breakthrough Waterc d c Rating from Scratch Phosphate force LC2 (N) based

Solventbased

K K + silicate seal K + sol-gel seal K + e-coat1 K + e-coat2 K + powder coat K + e-coat2 + powder coat Bare magnesium

10 9 10 10 10 10 10

9 9 9 10 10 9 10

0

0

6 10 10 10 10 10 10

14 ± 2 18 ± 1 18 ± 1 45 ± 2 N/T >100e >100e

4 4 4 4 4 4 4

5 5 5 5 5 5 5

N/T

N/T

N/T

N/T

Notes (N/T = not tested) a Keronite coating used was 10 μm for all samples. K = Keronite. b Tested according to ASTM B117-97 c Rating to ASTM D1654-92. 10 = no visible corrosion, 9 = <5% area corroded, 6 = 7-10% corrosion. d Rating to ASTM D1654-92. 10 = no corrosion creep, 9 = <0.5 mm peak creep. e Coating could not be removed with maximum load (100 N) on test machine f Rating to ASTM D3359-97. 5 = no paint removal, 4 = <5% removal.

558

6

Conclusions

1. The new Keronite process was applied to AZ91D plates with 5 μm/min coating speed. 2. Microscopic examination of the Keronite surface reveals a relatively smooth coating with fine pores, evenly distributed. 3. The average coating hardness was measured at 7.1 GPa for a 60 μm coating. 4. A 10 μm coating of Keronite improved the corrosion resistance of AZ91D. After 750 h salt spray testing, samples were rated 9 on scribed and 10 on unscribed samples. according to ASTM D1654 rating procedure. 5. The corrosion protection is enhanced by use of sealers or other top coats. 6. Keronite plus e-coat and/or powder coat was rated 10 even on scribed samples. 7. Inspection suggests that this combination will protect against galvanic corrosion also. 8. Keronite can allow magnesium components to be painted in a conventional steel-body paint process. 9. The scratch resistance of Keronite is known to be higher than other pre-treatments and can be enhanced 3 times with e-coat, and in excess of 9 times with powder coat. 10. A solvent-based paint was not removed from Keronite by tape test ASTM D3359-97.

6

References

[1] Shatrov, A., Samsonov, V., “Process and device for forming ceramic coatings on metals and alloys”, Patent Application PCT/GB2002/004305 [2] P. Shashkov, “Composite coatings on the base of Keronite ceramics on magnesium”, th 10 Magnesium Automotive and End-user Seminar, Sept 26-27, 2002 [3] J. Curran, Interim PhD report, Materials Science Laboratory, Cambridge University, Cambridge UK [4] Automotive e-coat and powder coat applied by Plastics Coatings Ltd, Dudley, UK [5] R. Beals, S LeBeau, O Roberto, S. Dewey, proc. SAE March, 2002 [6] Confidential correspondence with automotive vehicle maker [7] A. Yrokhin, Progress Report May-August 2002, University of Hull, UK

559

Corrosion Properties of PVD-coated Magnesium Die Cast Alloy AZ91D Part I: Investigation of the Corrosion Properties and the Corrosion Mechanisms of PVD Coated Magnesium Die Cast Alloy AZ91 for Tribological Applications H. Hoche 1, C. Blawert 2, E. Broszeit 1, C. Berger 1 1

Staatliche Materialprüfungsanstalt und Fachgebiet und Institut für Werkstoffkunde an der Technischen Universität Darmstadt, Darmstadt 2 GKSS Forschungszentrum Geesthacht GmbH, Geesthacht

Abstract Although the load bearing capacity of magnesium alloys is quite low in comparison to other construction materials, investigations of the authors showed good wear properties of PVD-CrN coated magnesium die cast alloy AZ91. In a certain parameter field the wear properties were in the same range as for PVD-CrN coated steel or titanium alloys [2]. Unfortunately the CrN coated specimens failed the corrosion tests due to crucial contact corrosion between the metalloid CrN and the less noble magnesium [3]. Developing PVD coating systems with better corrosion resistance leads to a duplex process, consisting of a plasma anodisation and an Al2O3 top coating which can both be performed in a modified commercial PVD unit. Even in the area of specific coating failures (Vickers indentations and scratches) the duplex coatings are not more affected by corrosion than the overall surface. Because the corrosion properties are not completely clarified yet, the aim of this work is to contribute to their scientific description. The following PVD-coating systems are included in the studies, which all meet the conditions of an acceptable wear resistance. • 9 μm CrN hard coating [1] • 3 μm TiN hard coating • 0.5 μm plasma anodisation layer & 3 μm Al2O3 coating [2] The publication is divided into two parts. Part I deals with the general coating and corrosion properties as well as with the electrochemical properties of the above mentioned coating systems. Some results are compared with a commercial available plasma-chemical anodisation produced in an aqueous electrolyte. The tested specimens were investigated by means of optical microscopy, SEM and EDX and with metallographic cross sections. Part II describes the results of a novel immersion technique to determine the time dependent corrosion behaviour of coated magnesium alloys. Furthermore contact corrosion properties of the above mentioned coatings according to DIN 50519 are presented. The specimens were tested in 3% NaCl solution and as contact materials the aluminium alloys AlMgSi 0,5 and AlSi9Cu3 and a roller bearing steel 100 Cr 6 were used. The tested specimens were investigated by means of optical microscopy, SEM and EDX.

560

1

Introduction

The care of brief resources and the economical use of primary energies are one of the great political and technological challenges of this new century in order to lower emissions and to reduce the increasing environmental pollution. Concerning this background, magnesium seems to be very prospective to fulfil these demands. Besides its high strength/ weight ratio, magnesium alloys combine the demands of good castability, good machiniability and a great recycling potential [4]. Because the resources of magnesium in the oceans and in the earth’s crust are nearly inexhaustible magnesium presents an ecological useful alternative to plastics. Because of the bad wear properties of magnesium alloys their application is limited up to now to the static component field. However, a considerably higher potential for the lasting fuel and emission reduction could be the application of magnesium alloys for moving parts by ensuring a sufficient wear resistance, e.g. with PVD hard coatings. Extensive research of the authors in deposition of reliable PVD hard coatings on magnesium alloys ([1], [2], [3]) showed comparable wear results in a certain parameter field in comparison to PVD coated titanium- and iron-based alloys. Considering the demands for the application of magnesium alloys in tribological systems the corrosion problematic becomes of crucial importance, particularly due to the PVD process related pinholes which can lead to crucial corrosion. If one thinks of decorative PVD coatings a reliable corrosion protection of the magnesium surface for applications in aggressive environments is indispensable if the coating is damaged. This occurs for example in automotive applications, when lambency water with abrasive particles such as loose chipping damages the coated surface. In this case the corrosion protection should not fail which leads to the need for developing coatings that do not loose their protective character even if the coating is partially damaged. A number of commercial deposition procedures are well established [6], [7]), which use aqueous electrolytes and are subjected to severe legal restrictions due to the use of harmful materials and the high amount of toxic waste. Furthermore the high roughness of these depositions complicates the subsequent PVD deposition or makes it rather impossible. The authors (IfW, Darmstadt) developed a prospective method to produce both corrosion and wear resistant coatings with the well established PVD technique. The process consists of two steps which can both be done in a slightly modified commercial PVD unit in one deposition run. The first step consists of producing a thin 0,5 μm thick plasma anodisation (PVD-PA) layer. The process is comparable with the anodisation in an aqueous electrolyte, substituting the aqueous electrolyte with an oxygen plasma. First, a positive BIAS voltage is applied to the magnesium substrate to be anodised ([1], [10]). In a following step a 3-4 μm Al2O3 hard coating is deposited by conventional RF sputtering. Within these studies the corrosion properties of the above mentioned coatings are compared to conventional hard coatings CrN and TiN.

2

Results

2.1

Coating Characteristics

The magnesium die cast alloy AZ91D was used as substrate material. The alloys of the AZtype (magnesium-aluminum-zinc) represent quantitatively the greatest group under the cast

561 alloys. Within this group AZ91D is the at this time most customary die casting alloy with a good combination of mechanical and physical properties and an outstanding castability. For the PVD surface treatments a commercial magnetron sputter unit type Alcatel SCM 601 was used. Target and substrates were connected to an r.f. power supply. In order to reduce the thermal load on the magnesium during deposition the substrate carrier was cooled. The samples were well prepared by grinding and polishing and afterwards they were cleaned in a two step ultrasonic bath with Isopropanol and n-Heptan. Prior to the deposition process a 20 minute sputter etching was applied to remove surface contaminations and to activate the surface for better coating adhesion and to improve the mechanical coating properties [3]. Table 1 shows the coating properties for the above mentioned coating systems and figure 1 shows SIMS depth profiles for a plasma anodised/Al2O3 coated and a CrN coated specimen. In order to improve the coating adhesion both the CrN and the TiN coating were deposited by sputtering a first layer of the pure metal and then gradually increasing the nitrogen flow. Therefore the correct terms for these coatings are Cr-CrN and Ti-TiN. In the case of the PA-Al2O3 the PVD plasma anodisation takes over the role of the primer. The nitrogen gradient and the PVD plasma anodisation layers can be seen in the SIMS depth profiles in figure 1. The Cr-CrN is electrically conductive, metallically bright and crystalline. The thickness of approx. 9 μm is in the upper coating thickness range of conventional PVDcoatings for wear protection. The deposition rate of Cr-CrN is in the range of 4.5 μm/h. The gold coloured Ti-TiN is also electrically conductive and crystalline. The coating thickness of approx. 3 μm is in the usual range of commercial TiN coatings due to the comparatively low deposition rates in the range of 1.5 to 2 μm/h. The adhesion of the TiTiN coatings is very weak and the coating detaches at values even lower than 1 N. In comparison, the PA-Al2O3 coating is electrically isolating, transparent and x-ray amorphous. With a thickness of 3 μm almost the upper limit for the deposition of ionic materials with the used deposition technique is reached. After optimising the partial pressures of argon and oxygen a maximum deposition rate of 1 μm/h was achieved. Increasing the deposition time decreases the deposition rates due to back sputtering effects so that after a deposition time of 8 h the coating thickness is only 4 μm. The 3 μm thick Al2O3 coating was deposited within 5 h. Leaving the small window of the optimal oxygen partial pressure rapidly halves the deposition rate. Although the roughness values are in the same range for all coating systems the roughness profile plots give evidence for the existence of pinholes in the Ti-TiN and the Cr-CrN coatings while the PA-Al2O3 plot shows a significant scattering which implied the existence of droplets onto the surface. These observations were validated with SEM investigations of the coatings surfaces. 2.2

General Corrosion Performance in the Salt Spray Test acc. to DIN 50021 SS

The general corrosion performances from the salt spray test for Cr-CrN and PA-Al2O3 coatings on AZ91 substrates were already demonstrated [2]. The Ti-TiN coatings show similar corrosion behaviour than the Cr-CrN coatings. For these tests pre damage was applied by producing a 10 mm scratch with a commercial scratch tester and then the samples were placed into the salt spray chamber. Figure 2 shows the surfaces after 96h salt spray testing. To get an impression of the salt spray test corrosion performance of the PVD-PA-Al2O3 coatings a conventionally anodised AZ91 sample (plasma chemical ano-

562 disation in an aqueous electrolyte) was included in the tests (figure 2 c). The thickness of the commercial anodisation is 25 μm and the coating porosity was sealed. Table 1: Properties of the Coating Systems.

adhesion roughness Ra Crystalline LcAE in N and Rz in μm

substrate-coating thickness in μm hardness system Gradient hard coat- HMplast in GPa or PA ing AZ91/ Cr/ CrN 2,2 7,0 41270 AZ91/ Ti/ TiN 0,7 2,3 42863 3,0 22078 AZ91/ PA / Al2O3 0,5 AZ91D

Al2O3

1E8

0,05 / 0,76 0,03 / 0,55 0,06 / 1,4

Yes Yes No

52

Cr

1E7

1E7

O Mg Al 64+66 Zn

100000 10000 16

O

24

Cr-CrN

1000000

AZ91D

14

N

27

24+25

Intensität

Al

1000

16

Plasmaanodisierung

27

1000000

Intensität

14 4 7

Mg Al Cr 64 Zn

100000

27 52

14

N

10000

27

Al

1000 100

64

Zn

24

Mg

10

100

64+66

Zn

24+25

Mg

10

1 0,0

0,5

1,0

1,5

2,0

2,5

3,0

3,5

4,0

4,5

0

5,0

1

2

3

4

5

6

7

8

9

10

11

12

13

14

15

Tiefe in μm

Tiefe in μm

b)

a)

Figure 1: SIMS depth profiles. a) AZ91/ plasma anodisation layer/ Al2O3, b) AZ91/ Cr/ CrN

10 mm

10 mm

a)

25 mm

b)

c)

Figure 2: Differently surface treated magnesium AZ91 samples, pre-damaged with scratch, after 96h salt spray testing. Samples not cleaned after unmounting from the test. a) 9 μm CrN, b) 0,5 μm PVD-PA + 2,7 μm Al2O3, c) conventional plasma-chemical anodisation, thickness approx. 25 μm

563 0,5 μm PVD PA + 2,7 μm PVD Al2O3

25 μm conventional plasma chemical anodisation

Cross section A-A, Depth 3 mm, Load 30 N

A-A

A-A Cross section B-B, Depth 5 mm, Load 50 N

B-B

B-B

Cross section C-C, Depth 7 mm, Load 70 N

C-C

D-D Cross section D-D, Depth 9 mm, Load 90 N

C-C

D-D

Figure 3: SEM figures of the scratches after 96h salt spray testing and cross sections through the scratches. Left hand: PVD plasma anodisation and Al2O3 hard coating, right hand: conventional plasma-electrolytical oxidation.

While the initial surfaces of the PVD-PA-Al2O3 and the commercial anodisation in figures 2 b) and c) are widely intact, the surface of the Cr-CrN coated specimen is charged with thick foam like corrosion products (figure 2 a). Optical surface inspection of the as tested surfaces shows a better corrosion performance of the PVD-PA-Al2O3 coated sample in comparison to the plasma electrolytic anodisation. A detailed examination of the corrosion attack is done with SEM and with metallographic cross sections (figure 3). The SEM-investigations show a higher amount of corrosion products on the scratched surface of the conventional plasma chemical anodisation. This would be expected because

564 of the much higher coating thickness. Up to a scratch depth of 3 mm the corrosion attack from the cross sections is negligible for both the PVD-PA-Al2O3 and the conventional plasma chemical anodisation. The cross section at 5 mm depth shows one pitting like corrosion attack at the PVD treated surface while the structure of the conventional anodisation shows a overlap like structure due to the plastic deformation of the substrate. Both samples show a distinct amount of corrosion starting at a scratch depth of 7 mm. The corrosion attack of the PVD treated surface is mainly homogenous and moderate through the whole cross section of the scratch. On the other hand the conventional product shows crucial cauliflower-like corrosion attack with diameters in the range of 100 μm. Increasing the load leads to the formation of cracks inside the structure of the substrate by exceeding the plasticity of the substrate. The main corrosion attack at the conventional plasma chemical anodisation is located in the areas of maximum tensile stresses due to the plastic deformation of the scratches. In the compressive stress areas in the ground of the scratch where plastic deformation of the substrate is at its maximum, there is only less corrosion. For the PVD-PA-Al2O3 coating this dependence is not valid because the corrosion is uniform. One explanation can be a uniform cracking of the “thin” PVD coating and as a result of this somehow maintaining a corrosion protection effect while the “thick” conventional plasma anodisation prefers cracking in the areas of maximum tensile stresses where the high corrosion attack was observed. As known from previous salt spray tests of Cr-CrN coated AZ91 specimens the surface is crucially damaged which implies that the initial scratch is also crucially damaged. After 6+ removing the corrosion products by etching in Cr -acid, the surface is – as expected – crucially affected by corrosion except for the scratch itself. This behaviour was observed with all tested CrN based coatings (figure 4a). Looking at the etched cross section there is significant plastic deformation of the microstructure up to a depth of 30 μm below the scratched surface (figure 4b). This seems to lead to a densification of the network of the intermetallic Mg17Al12 Phase and therefore seems to affect the corrosion properties in a positive way. In further investigations it should be checked, whether a rolling of the surface can improve the corrosion properties of PVD-coated magnesium alloys.

a)

b)

Figure. 4: a) Plastic deformation of the microstructure due to the scratch. b) crucial corrosion next to the scratch at the CrN coated AZ91, no remarkable corrosion of the scratch.

2.3

Electrochemical Behaviour

Figure 5 shows the current vs. potential curves of PVD-PA-Al2O3, TiN and CrN coated AZ91 specimens in 5% NaCl solution. Using an exposed surface area of 1.23 cm2 and the Tafel-equation a corrosion rate can be determined from the cathodic branch of the curves.

565 The calculated corrosion rates are included in figure 5. The calculated corrosion rate for the PVD-PA-Al2O3 is 0.02 mm/year and therefore a factor between 100 and 2000 lower than for the uncoated alloy (2.2 mm/year) and the Ti-TiN (13.5 mm/year) and the Cr-CrN (37.7 mm/year) coatings. The anodic branches which represent the active dissolution are comparable for all samples shown in figure 5 and no differentiation is possible. Reducing the NaCl concentration to 0.1% and 0.01% leads to comparable results with a nearly similar anodic behaviour of the tested specimens. In order to gather knowledge about the corrosion initiation additional measurements were carried out in 0,1% and 0,01% NaCl and aborted at a current density of 50 μA/cm². Afterwards the surfaces were investigated by optical microscopy, SEM and metallographic cross sections. The results form the tests in 0,1% NaCl are presented in figure 6. Although the run of the curves are comparable, there are crucial differences between the different coating systems. The NaCl concentration does not affect the macroscopic corrosion appearance. As one can clearly see the macroscopic inspection shows no corrosion attack on the PVD-PA-Al2O3 coating but remarkable corrosion on the Cr-CrN and the Ti-TiN coated surfaces, where the attack of the Ti-TiN coated sample is lower than on the Cr-CrN coated sample. The SEM figures clearly show that there is detachment of the Cr-CrN and the TiTiN coating at the corrosion pits. This assumption is confirmed by the cross sections, where electrolyte penetrates the interface at growth defects leading to corrosion in the interface. The larger volume of the corrosion products leads to a coating lift of. There is interface corrosion over the whole cross section area at the Ti-TiN and the Cr-CrN. In contrast, the interface of the PVD-PA-Al2O3 is still intact and only one spot shows corrosion products on the coatings surface. The SEM figure of the PVD-PA-Al2O3 also shows a growth defect in the form of a droplet which is unavoidable for this coating type. The defect size on the PVD-PA-Al2O3 is below 15 μm while the defect size on the CrCrN is > 200 μm and on the Ti-TiN > 15 μm. An estimation of the absolute defect number on the tested surfaces leads to a ratio of PVD-PA-Al2O3 : Cr-CrN ≈ 4 : 1. Considering the defect sizes leads to an absolute defect area ratio of PA-Al2O3 : Cr-CrN ≈ 1 : 100. Ignoring the influence of current density of the intact coating area the current density at the defects of the PVD-PA-Al2O3 coated sample is hundred times larger than for the CrN although the current densities related to the whole testing areas are in the same ranges at the applied potential. 100

10

current in mA

1

Cr-CrN 37.7 mm/year Ti-TiN 13.5 mm/year AZ91D 2.2 mm/year

0,1

0,01

PA-Al2O3 0.02 mm/year

0,001

0,0001 -1600

-1500

-1400

-1300

-1200

potential in mV

Figure. 5: Current- vs. Potential curves.

-1100

-1000

566 Cr-CrN

Ti-TiN

2,5 cm

Cross section

SEM

Macroscopic Inspection

PVD-PA + Al2O3

encapsulation PA+Al2O3 AZ91

Fig. 6: Characterization of the samples after testing in 0,1%-NaCl solution, aborted at a current density of 50 μA/cm²

Literature [1] H. Hoche, H.J. Schroeder, H. Scheerer, E. Broszeit, C. Berger, “Tribological Studies of CrN coated Magnesium AZ91 at Temperatures up to 250 °C”, Advanced Engineering Materials 4 (2002), 42-51 [2] H. Hoche, H. Scheerer, E. Broszeit, C. Berger, „Erschließung von Mg-Legierungen für tribologisch-korrosive Komplexbeanspruchung mittels durch Plasmaanodisation erzeugter Korrosionsschutzschichten“, Mat.-Wiss. u. Werkstofftech., 33 (2002), 395-403 [3] H. Hoche, H. Scheerer, R. Fritsche, A. Thißen, S. Flege, E. Broszeit, C. Berger, H.M. Ortner, W. Jaegermann, „Untersuchung des Sputter-Ätzens auf die Eigenschaften von PVD-CrN Hartstoffschichten auf Magnesium AZ91hp“, Mat.-Wiss. und Werkstofftech, 33 (2002), 132-141 [4] Aghion, E., Bronfin, B., „Magnesium Alloys Development towards the 21st Century“, Mat. Sci. Forum Vold. 350-351 (2000), S. 19-28 [5] Haferkamp, H., „Magnesiumkorrosion – Prozesse, Schutz von Anode und Kathode“ in: Moderne Beschichtungsverfahren, S. 242-260 [6] Thoma, M., „Oberflächenschutz für Magnesiumwerkstoffe zur Verbesserung des Korrosions- und Verschleißverhaltens“, Metalloberfläche 38 (1984) 9, p. 393-99

567 [7] Jelinek, T.W., „Fortschritte in der Galvanotechnik – Eine Auswertung der Fachliteratur 2001/2002“, Galvanotechnik 1/2003 (2003), S. 46-74 [8] Kuhn, A., “Plasmaanodisation von Magnesiumlegierungen”, Galvanotechnik 94 (2003), S. 1114-1122 [9] D’Angelo, S, AIFM galvanotecnica 12 (2001), S. 37-41 [10] G.J. Tibol, R.W. Hull, “Plasma Anodized Aluminium Oxide Films”, J. Electroch. Soc. (1964), 1368-1372

568

Application of Magnesium Alloys Coated with High Purity Mg for Engineering Usage A. Yamamoto*, R. Tamai** and H. Tsubakino* * Graduate School of Engineering, Himeji Institute of Technology, Himeji, Japan ** Graduate Student of Himeji Institute of Technology, Himeji, Japan

1

Introduction

The present authors have reported on a technique of deposition coating on magnesium alloys with high purity magnesium for improving the poor corrosion resistance in magnesium alloys [1-5]. It was shown about sixty years ago by Hanawalt et al. [6] that heavy metal impurities deterioratively affect the corrosion resistance in magnesium and its alloys. Vacuum evaporation method utilized in our technique intrinsically includes a purification process, which effectively decrease concentrations of heavy metals. Salt immersion tests on the coated specimens by our technique and un-coated magnesium alloys showed that corrosion rate on the coated specimen was less than that on AZ91E magnesium alloy and was comparable to that in high purity magnesium (6N-grade) [5]. Increase in corrosion resistance in the coated specimen has been explained also by the authors that local corrosion, the filiform corrosion [7], occurs in the un-coated AZ31 alloy and 3N-grade pure magnesium, while the filiform corrosion has not been observed in 6N-Mg and the coated magnesium alloy [3]. As a result, rough surfaces are formed in the former, although smooth surfaces are formed in the latter by corrosion reaction, which lead to poor or superior corrosion resistance in the specimens, respectively [4]. The filiform corrosion occurs near inclusions and tends to proceed along grain boundaries [5], that is, inhomogeneity in microstructures and chemical compositions such as grain boundary segregations provides favorable sites for the filiform corrosion. In the coated magnesium film, there is no inhomogeneity in chemical compositions, because it is composed of high purity magnesium, so the corrosion resistance can be improved. Additionally, the following results have been obtained on the deposition coated AZ31 alloy: hot-press treatment of the coated specimens increased corrosion resistance of the specimens by decreasing micro-structural defects within coated films [8], and when a solid recycling process, in which scraps were hot-extruded, was applied to the deposition coated specimens which was solid recycling process using a hot extrusion method was applied to the deposition coated AZ31 alloy, the recycled specimen showed almost the same tensile properties as those in the virgin materials [9]. Fundamental conditions for deposition coating technique are seemed to be established by the above mentioned studies. However, the specimens used were too small, about 3 10 x 10 x 5 to 60 x 40 x 5 mm , to apply to engineering products. Therefore, we developed a new deposition coating apparatus which enables one to use a large substrate with about 250 x 250 mm2 in area for coating. In the present study, preliminary results obtained by using the new apparatus are shown.

569

2

Experimental procedures

Schematic illustration for the newly developed apparatus for deposition coating of high purity magnesium on magnesium alloys are shown in Fig. 1. There are two vacuum chambers for evaporation and for deposition, which can be evacuated about 10-1 and 10-3 to 10-4 Pa, respectively. There is a gate valve between them, which is used to exchange a substrate or an evaporation source independently. A substrate is attached on a setting plate made with stainless steel plate which is heated at the temperature range between R.T. to 673 K. The evaporation source is set in a graphite crucible which is heated up to 873 K. In order to avoid deposition of magnesium on a wall in the deposition chamber except the substrate plate, a cylindrical shield wall is set around the side wall in the deposition chamber and is heated at higher temperature than that of the substrate. The substrate with about 250 x 250 mm2 in area can be deposition coated. Substrate heater Substrate (Mg alloy)

Deposited Mg film

Deposition chamber: 10-3-10-4 Pa Substrate temperature: <673 K Deposition area of substrate: 250 x 250 mm2

Shield wall Heater for shield wall

Evaporation chamber: 10-1 Pa Crucible temperature: <873 K Weight of evaporation source: 40 g

Deposition chamber

Gate valve

DP

Evaporation source (Mg)

RP

Crucible heater Graphite crucible

MB

RP

(MB: Mechanical booster pump)

Evaporation chamber

Figure 1. Schematic illustration of the newly deposition coating furnace.

Appearance of the apparatus is shown in Fig. 2 (a), and the deposition and evaporation chambers are shown in Fig. 2 (b) and (c), respectively. For the substrate, AZ31 magnesium alloy was used, while 3N-Mg was used for evaporation source, chemical compositions of which are listed in Table 1. Table 1. Chemical compositions of the AZ31 alloy and 3N-Mg (mass%).

3N-Mg AZ31

Al

Zn

Mn

Si

Cu

Ni

Fe

0.002 2.96

0.004 0.828

0.005 0.433

0.001 0.004

0.002 0.004

<0.001 <0.001

<0.001 0.002

570 (b) Deposition chamber (a) Substrate Deposition chamber (c) Control unit Evaporation chamber Vacuum system

Evaporation chamber

Figure 2. Newly developed deposition coating furnace. (a) Appearance of the apparatus, (b) deposition chamber and substrate plate, (c) evaporation chamber.

The setting plate for substrates is shown in Fig. 3, which was used for the deposition experiment carried out at the conditions with the temperatures of substrate and evaporation source were 427 and 823 K, respectively. There are many small holes for setting the substrates with various sizes. In this case, the substrates of AZ31 sheet with 3 about 20 x 30 x 0.8 mm were used, the traces of the substrates can be seen as indicated by the arrows. Substrates with the above mentioned one, 50 x 50 x 0.8 mm3 and 250 x 250 x 0.8 mm3, and also a die-cast part made with AZ91 alloy for a cell phone were attempted to be deposition coated.

50 mm

Figure 3. Setting plate made with stainless steel sheet after being used for deposition coating experiment on the substra3 tes of AZ31 alloy with 20 x 30 x 0.8 mm in sizes. The arrows show the traces of the substrates.

571

3

Results and discussion

In the previous paper [5], the optimum temperature range for the substrate was reported as about 500 to 620 K, while in the case of the present new furnace, appropriate temperatures for deposition were in the range of about 323 to 523 K. Deposited film was formed below 323 K, but adhesion between the film and the substrate was poor. On the other hand, evaporation occurred at higher temperatures than 523 K. Microstructures of the deposition coated surfaces prepared at the substrate temeperatures of 323, 373 and 473 K are shown in Fig. 4. Temperature of the evaporation source was 823 K. Deposition time, which means a period of opening the gate valve, was 1.8 ks. Facetted flake-like microstructures are observed on the specimen prepared at 323 K (a), the size of the flakes increased with increasing temperature (b), and the flake-like microstructure was seemed to be changed into relatively continuous characteristics with small voids at 473 K (c). The small voids, however, were observed only on the surface region and were not formed penetrating the deposited layer as shown in the cross section (d). The temperature for the evaporation source is also lower than that in the previous study, 980 K [5]. In both the cases, 3N-Mg is not melted but sublimation occurs depending on the pressure in the furnace. The differences in the optimum temperatures for evaporation

(a)

(b)

2 μm (c)

2 μm (d)

Deposited layer Substrate 2 μm

5 μm

Figure 4. Microstructures of the deposition coated surfaces fabricated at the temperatures of 323, 373 and 473 K, (a), (b) and (c), respectively, and the cross section of the deposited surface coated at 473 K, (d).

572 and deposition are considered to be due to the differences in degree of vacuum. Although the measured values in both the cases are about the same, 1 x 10-3 Pa, the position for a vacuum gauge was relatively apart from the furnace in the case of the previous study, in comparison with the present study. Lowering the pressure in the furnace means lowering a partial pressure of oxygen and would prevent oxidation of the substrate surface, which beneficially affects the microstructure of the interface between the deposited layer and substrate. In Fig. 4 (d), it is difficult to discern the interface, which is seemed to be comparable to that in the HIP treated specimen [8]. And also the defects such as penetrating voids through the deposited film are not formed in the present case, which ensures improving corrosion resistance because anodic corrosion in the deposited film near the penetrating voids would be avoided. The main advantages for applying magnesium alloys to portable electronic devices are their lightness, high thermal conductivity and recycle ability, in comparison with those in polymer products. Figure 5 (a) shows such an example, which is a die-cast product for an inner part of a cell phone. It is obvious that flow marks are formed along paths of molten alloy, which is a typical defect in die-casting and it is often difficult to cover with surface treatments such as anodic treatments and paintings. Deposition coating was attempted to apply to the part in order to cover the flow marks. Coating was carried out at the temperatures of 473 and 843 K for the specimen and evaporation source, respectively, for 1.2 ks. Coating was successfully completed and the resultant coated specimen is shown in Fig. 5 (b), in which the flow marks are covered and not be observed. Cross section of the coated specimen is shown in Fig. 6. Thickness of the coated layer is about 2.5 μm. There are no defects at the interface between the coated layer and the substrate as like as that shown in Fig. 4 (d). The surface of the substrate used for Fig. 4 was finished by mechanically polishing with emery paper, #2000, while the surface of the die-cast sample was only rinsed with ethanol for degreasing. The deposition coating technique can be applied as an alternative process for primer coating. (a)

(b)

20 mm Figure 5. Die cast parts for a cell phone: as cast and rinsed (a), and deposition coated (b).

573

Coated layer Interface

Die-cast sample

5 μm

Figure 6. Cross section of the deposition coated die-cast product shown in Figure 5 (b).

4

Summary

Newly developed deposition coating apparatus has been introduced, which enables one to carry out coating on 250 x 250 mm2 in area. Defects between the coated layer and the substrate have been decreased, which is considered to be due to high vacuum in the coating chamber. As an example for applying this technique to engineering products, a diecast part for a cell phone has been attempted to be coated, which resulted in forming a deposited layer without any defects between the layer and the substrate. The deposition coating technique would be a promising process for primer coating instead of a conventional one as well as a surface finishing.

5

References

[1] Yamamoto, A.; Watanabe, A.; Sugahara, K.; Tsubakino, H.; Fukumoto, S., Scr. Mater. 2001, 44, 1039-1042. [2] Yamamoto, A.; Watanabe, A.; Sugahara, K.; Fukumoto, S.; Tsubakino, H., Mater. Trans. 2001, 42, 1237-1242. [3] Yamamoto, A.; Watanabe, A.; Sugahara, K.; Fukumoto, S.; Tsubakino, H., Mater. Trans. 2001, 42, 1243-1248. [4] Tsubakino, H.; Yamamoto, A.; Sugahara, K.; Fukumoto, S.; Mater. Sci. Forum 2003, 419-422, 915-920. [5] Tsubakino, H.; Yamamoto, A. Fukumoto, S.; Watanabe, A.; Sugahara, K.; Inoue, H., Mater. Trans., 2003, 44, 504-510. [6] Hanawalt, J. D.; Nelson, C. E.; Peloube, J. A., Trans. Metall. AIME., 1942, 147, 273299. [7] Ghali, E. Mater. Sci. Forum, 2000, 350-351, 261-272. [8] Fukumoto, S.; Sugahara, K.; Yamamoto, A.; Tsubakino, H., Mater. Trans., 2003, 44, 518-523. [9] Chino, Y.; Yamamoto, A.; Iwasaki, H.; Matuchi, M.; Tsubakino, H., Mater. Trans., 2003, 44, 578-582.

574

Corrosion and Fatigue Behaviour of the Magnesium Die-Cast Alloy AZ91 hp after Surface Treatment C. Dindorf, C. Müller Technische Universität Darmstadt, Germany

1

Introduction

Surface treatment by shot-peening (SP) has been known for many years as an established process for increasing the fatigue life of steels and non-ferrous alloys [1, 2]. The impact of the shot media on the surface leads to plastic deformation and compressive residual stresses in the surface layer, obstructing the nucleation and propagation of fatigue cracks in or near the surface [3-5]. The effect of SP on lightweight alloys is of particular interest in order to extend their field of application as mechanically loaded components in the automotive and aerospace industry. However, the fatigue properties of shot-peened magnesium alloys have not yet been intensively covered in the literature. Main contributions on that subject can be found in [6-9]. In the SP process small balls, usually made out of steel, are accelerated towards the surface. As a consequence, residual iron particles remain on the surface after the treatment. With an electrolytic solution present, severe galvanic and pitting corrosion sets in [10]. According to [11] this detrimental effect of SP concerning corrosion can be avoided by (1) peening with a chemically inert material such as glass or ceramics, (2) chemical removal of the iron particles, or (3) secondary peening with glass beads or grinding particles made of alumina. The aim of the present study is to apply the latter method and investigate the fatigue and corrosive behavior after secondary peening with glass beads.

2

Experimental Procedure

The material employed for the investigations is the magnesium die-cast alloy AZ91 hp. The material was cast as cylindrical fatigue specimens that have a gauge diameter of 6 mm and a test length of 85 mm (Fig. 1).

Figure 1: Die-cast fatigue specimens

575 The specimens were fatigued at 50 Hz in a rotating bending beam setup, thus with a stress ratio of R = -1. Where stated, specimens were precorroded for three hours in an immersion test with a 5% NaCl aqueous solution before fatigue testing. The SP treatment was performed at the University of Technology Cottbus, using ferritic shot SCCW 14 (ball diameter: 0.40 mm) with an Almen intensity of 0.41 mmN. For secondary peening glass beads have been employed, having a diameter of 0.07 to 0.10 mm, subsequent to the aforementioned SP treatment with ferritic shot.

3

Results and Discussion

3.1

Surface Treatment

Looking at SEM-micrographs of the surfaces of the three kinds of specimens investigated (Fig. 2), there are clear differences in the surface quality. Fig. 2a shows the as-cast on the left and the shot-peened part of a specimen on the right. In the as-cast area there is hardly any contrast due to its smoothness. The shot-peened material (Fig. 2a,b) shows a much rougher, crater-like surface, caused by the impact of the steel balls. Compared to that, the surface after SP and secondary peening is less jagged and considerably smoother (Fig. 2c). The much smaller and less hard glass beads appear to level out the cratered surface caused by SP with ferritic shot.

Figure 2: SEM-Micrographs of shot-peened surfaces. (a) transitory area between the as-cast grip section and the shot-peened test length; (b) with ferritic shot; (c) with ferritic shot plus secondary peening with glass beads

The results of an EDX analysis (Table 1) show that the as-cast surface has no detectable quantity of iron on it. After SP the iron concentration is significant, by more than two orders of magnitude higher than in the normal composition. The secondary peening with glass beads results in a cleansing of the surface, i.e. the amount of iron is significantly reduced to one fourth of the value after SP. Table 1. EDX analysis of the different specimen surfaces

Element / wt.% Fe Si

As-cast 0,0 0,1

SP 0,8 0,3

SP + glass beads 0,2 0,8

During the secondary peening process the surface is covered with silicon residue from the glass beads. For the as-cast specimen the concentration of Si – the primary element in the glass beads – is within limits typical of AZ91 hp. After secondary peening the Si concentration is highly increased compared to the as-cast concentration.

576 Vickers microhardness profiles of shot-peened and secondary peened specimens have been produced (Fig. 3). For both peening conditions the hardness is significantly increased in the outer 200 μm. After SP the rise in surface hardness is around 80% whereas there is an increase of 90 to 120% for secondary peened samples. Thus, the secondary peening seems to have an additional strain hardening effect. The high variation of the hardness between the outer rims can be ascribed to the presence of casting pores.

Figure 3: Microhardness profile along the diameter of the shot-peened specimens (a) with ferritic shot; (b) with ferritic shot plus secondary peening with glass beads

3.2

Corrosion

Figure 4 depicts the corrosive effect of a 3 hour long immersion test in a 5% NaCl aqueous solution on the three different kinds of specimens investigated in this study. The as-cast specimen does not show any significant corrosive attack on the surface (Fig. 4a). There are only some small spots scattered across the upper half of the picture. By and large, though, the surface appears to be unaffected by the immersion test. The SP, however, leads to severe pitting in a corrosive environment (Fig. 4b). By SP the die skin of the cast specimen is removed. In previous studies, the die skin was proven to have a beneficial effect on the corrosive behavior, acting as a protective layer [12]. During the immersion test the solution resembles sparkling water due to the high rate of gas bubbles rising from the shot-peened surface.

Figure 4: Micrographs of specimen surfaces corroded for 3 hours in 5% NaCl solution (a) as-cast; (b) SP with ferritic shot; (c) SP with ferritic shot plus secondary peening with glass beads

The effect of pitting appears to be almost as pronounced after secondary peening with glass shot (Fig. 4c). The number, width and depth of the pits are smaller, though. Yet, compared to the primarily uniform corrosion of as-cast samples the corrosive attack is still

577 considerable. Thus, even though the iron concentration has been reduced to one fourth by secondary peening, the removal of contamination is not sufficient to suppress pitting corrosion and its consequences. Apart from a comparative analysis of micrographs of the corroded surfaces the corrosive attack by pitting or galvanic corrosion cannot be quantified. The usual characterization of corrosive behavior through the means of current-density-potential curves is not applicable in this case as only the external current, and not the local one, can be measured. However, this method was successfully employed to analyze the amount of surface contamination through the SP process. The curve of the sample after SP is shifted towards more noble potentials compared to the curve of the as-cast specimen. The secondary peened material runs with its curve inbetween the two. The shift is caused by the amount and standard potential of surface contamination. So, the position of the curve for the secondary peened specimen shows the partial removal of the relatively noble iron. 3.3

Fatigue Properties

The fatigue properties for the uncorroded and peened specimens are more or less identical. Even though secondary peened specimens appear to have a longer life-time for higher loads both kinds of uncorroded specimens share the same endurance limit of approximately 100 MPa (Fig. 5). Thus, the potentially higher strain hardening in the secondary peened surface layer, as discussed above, does not seem to have a significant influence on the fatigue properties. The smoothening of the surface after secondary peening (Fig. 2b,c) has apparently no major impact on fatigue, either. The fatigue resistance of the corroded specimens is drastically decreased. Again, secondary peened samples show a slightly superior fatigue life-time for higher loads, yet the endurance limit appears to be the same for both corroded peening conditions, lying below an unacceptable value of 30 MPa. The endurance limit could not be determined as with the setup employed for this study the accurate application of such a low load could not be warranted.

Figure 5: Fatigue properties of the different kind of specimens investigated

578 The virtually identical fatigue properties of the precorroded specimens after SP as well as secondary peening show that the removal of iron contamination by secondary peening was not sufficient, and that the decreased density of corrosion pits for secondary peened samples (Fig. 4b,c) does not improve fatigue properties significantly in comparison to corroded samples after normal SP.

4

Summary and Conclusions

A surface treatment in the form of shot-peening (SP) was performed with ferritic shot on specimens of the die-cast magnesium alloy AZ91 hp in order to increase the fatigue life of the material. The SP treatment leads to severe pitting corrosion in chloride immersion tests, therefore a method for removing the detrimental iron particles was investigated while trying to maintain the improved fatigue properties, if possible. The results led to the following conclusions: • Through secondary peening with glass beads the iron contamination from SP can be reduced by approximately 75%. The surface appears to be smoothened out compared to the SP condition. • The removal of iron from the surface is, however, not sufficient to prevent pitting corrosion in an immersion test. • The resulting fatigue properties are not significantly superior to those of specimens after SP, despite a lower density of corrosion pits. • For both precorroded conditions the endurance limit is drastically decreased by at least 70% compared to uncorroded specimens after SP. • Contradictory to [11], secondary peening with glass beads can be ruled out as a means to prevent subsequent pitting corrosion as well as unacceptable fatigue properties. Further studies will deal with the effect of pickling the specimens after SP. Subsequent immersion and fatigue experiments will be performed and evaluated.

5

References

[1] Cary, P.E., in Conference Proceedings: ICSP-1, www.shotpeening.org 1981, p. 2328. [2] Niku-Lari, A., in Conference Proceedings: ICSP-1, www.shotpeening.org, 1981, p. 237-247. [3] Benedetti, M.; Fontanari, V.; Höhn, B.-R.; Oster, P.; Tobie, T., International Journal of Fatigue 2002, 24, 1127-1136. [4] Curtis, S.; de los Rios, E.R.; Rodopoulos, C.A.; Levers, A., International Journal of Fatigue 2003, 25, 59-66. [5] Drechsler, A.; Dörr, T.; Wagner, L., Materials Science and Engineering 1998, A243, 217-220. [6] Wagner, L., Materials Science and Engineering 1999, A263, 210-216. [7] Dörr, T.; Hilpert, M.; Beckmerhagen, P.; Kiefer, A.; Wagner, L., in Conference Proceedings: ICSP-7, www.shotpeening.org 1999, p. 153-160.

579 [8] Altenberger, I.; Scholtes, B., in Surface Treatment IV: Computer methods and Experimental Measurements (Ed.: C.A. Brebbia), WIT Press, Southampton (UK), 1999, p. 281-290. [9] Noster, U.; Altenberger, I.; Scholtes, B., in Surface Treatment V: Computer methods and Experimental Measurements (Ed.: C.A. Brebbia), WIT Press, Southampton (UK), 2001, p. 3-12. [10] Müller, C.; Rodríguez, R., in Shot-Peening (Ed.: L. Wagner), Wiley-VCH, Weinheim, 2003, p. 271-278. [11] Hammersley, G.; Hackel, L.A.; Harris, F., Optics and Lasers in Engineering 2000, 34, 327-337. [12] Müller, C.; Deinzer, G.H.; Koch, R., in Magnesium Alloys and their Application (Ed.: K.U. Kainer), Wiley-VCH, Weinheim, 2000, p. 457-462.

580

Corrosion Behaviour in Artificially Corrosion-Oxidation treated Mg Alloys A. Yamamoto*, H. Inaoka** and H. Tsubakino* * Graduate School of Engineering, Himeji Institute of Technology, Himeji, Japan ** Graduate Student of Himeji Institute of Technology, Himeji, Japan

1

Introduction

In order to improve the corrosion resistance in magnesium and its alloys, the present authors have proposed new techniques for surface treatments. The first one is the deposition coating technique [1], in which surfaces of magnesium alloys are covered with high purity magnesium by applying vacuum deposition method which intrinsically includes purification process of magnesium resulting in improving corrosion resistance. The second one is a corrosion-oxidation technique [2, 3], in which specimens of magnesium alloys are immersed into aqueous solutions with relatively high pH, and then the specimens are heated in air. Magnesium hydroxide film is formed on the surface of the specimen by salt immersion, which is changed into magnesium oxide by heating to cover the surface. Advantages in the corrosion-oxidation technique are intrinsically included in the method, that is, only the aqueous solution and air are used in this technique, which enables one to cover surfaces in narrow gaps, internal surfaces of complicated-shaped products and so on without any directionality, and to apply easily the technique to large size products. Additionally, it should be mentioned that the technique is adapted to concepts for environmental protection and energy saving, harmful regents are not used and conventional furnaces are only used. In the present study, microstructures formed by the corrosion-oxidation technique are mainly investigated on AZ31 magnesium alloy as well as 3N-Mg, AZ61 and AZ91E alloys.

2

Experimental procedures

Chemical compositions of the alloys used are listed in Table 1. Specimens with about 10 x 10 x 5 mm3 in sizes were mechanically polished using emery papers and finished with #2000, and chemically polished with phosphoric acid : ethanol = 3.5 solution. The specimens were immersed into an aqueous solution with (10% NaCl + 10% NaOH), pH of which was about 12, the immersion was carried out at R. T. for 3.6 ks. After immersion, the specimens were rinsed and dried, and then heated in air at 673 K for 3.6 ks. Surfaces of the specimens were observed by a scanning electron microscope (SEM: Hithchi, S-900) attached with an energy dispersive X-ray analyzer (EDX). Products due to corrosion and oxidation were observed by a transmission electron microscope (TEM: JEOL, JEM-200CX) under an accelerating voltage of 200 kV. Thin foils for TEM observations were prepared by using a microtome, and also by the scratch method reported in the previous paper [2].

581 Corrosion resistance in the specimens were evaluated by in-situ laser microscopic observations. That is, specimens are immersed into 1% NaCl solution under observation by means of a laser microscope (Lasertec, 1LM15), and corrosion behavior is observed through a layer of the solution with about 3 to 5 mm in depth. Corrosion reactions for magnesium are represented as follows [4]: Mg  Mg2+ + 2e2H2O + 2e-  H2 + 2(OH)Mg2+ + 2(OH)-  Mg(OH)2

(1) (2) (3)

Evolution of hydrogen bubbles due to the reaction (2) can be observed, while other reactions are not observed. Table 1. Chemical compositions of the AZ31 alloy and 3N-Mg (mass%).

3N-Mg AZ31 AZ61 AZ91E

3

Al 0.005 2.83 6.26 9.18

Zn 0.004 0.93 0.68 0.66

Mn 0.003 0.003 0.26 0.22

Si 0.005 0.007 0.020 0.019

Cu <0.001 0.003 0.002 0.006

Ni <0.001 <0.001 <0.001 <0.001

Fe 0.002 0.004 0.004 0.002

Results and discussion

Microstructures of AZ31 alloy after electropolishing, immersion in (10% NaCl + 10% NaOH) solution with about pH 12 at R. T. for 3.6 ks and then heated in air at 673 K for 3.6 ks are shown in Fig. 1 (a), (b) and (c), respectively. Corrosion products were formed on the surface by salt immersion (b) but they disappeared by the following heat treatment as shown in (c). Dark small contrasts in (c) are confirmed to be due to magnesium oxide by EDX analyses. It is worth to notice that grain boundaries in the specimen can be seen (a)

(b)

(c)

100 μm Figure 1. Microstructures in AZ31 alloy. After polished and etched (a), corroded in (10% NaCl + 10% NaOH solution) (b), and then heated in air at 673 K for 3.6 ks (c).

582 even after oxidation treatment (c), while in the case of a directly oxidized specimen, surfaces of the specimen were covered with thick oxide layers and details of the microstructures were not observed. Excess oxidation reaction seems to be suppressed by the corrosion treatment. Microstructures on surfaces of 3N-Mg, AZ31, AZ61 and AZ91E magnesium alloys after corrosion treatment and also corrosion-oxidation treatment are shown in Fig. 2 (a) (d) and (e) - (h), respectively. Flake-like microstructures are observed in all the corrosiontreated specimens, which were not significantly changed by oxidation treatment, except in AZ31 alloy (f), but in other view fields the microstructure was observed to be not so changed. (a)

(b)

(c)

(d)

(e)

(f)

(g)

(h)

0.1 μm

Figure 2. Microstructures of surfaces on 3N-Mg, AZ31, AZ61 and AZ91E magnesium alloys after corroded in (10% NaCl + 10% NaOH) solution, (a), (b), (c) and (d), respectively, and then heated in air at 673 K for 3.6 ks, (e), (f), (g) and (h), respectively.

Results of in-situ observations of corrosion behaviors on un-treated (only electropolished) and corrosion-oxidation-treated specimens of AZ31 alloy are shown in Fig. 3. In the case of the un-treated specimen, hydrogen bubbles were immediately evolved in 1% NaCl solution on the surface of the specimen as shown in (a), while in the case of the corrosion-oxidation treated specimen, evolution of bubbles were suppressed and only several bubbles were observed on the surface with about 10 x 10 mm2 in area for about 3.6 to 7.2 ks. After 900 s, on the surface of the un-treated specimen, active corrosion suddenly occurred as shown in Fig. 2 (c), in which positions for evolution of bubbles advanced remaining dark inactive tails, those are the characteristics of the filiform corrosion [5, 6]. On the other hand, the time for starting the filiform corrosion was prolonged in the corrosionoxidation treated specimen, 18 ks in the case of the specimen shown in (d). It has been shown in the previous paper that heads of the filiform corrosion advance along grain boundaries in the un-treated AZ31 alloy [6], and that grain boundary segregation is con-

583 sidered to provide a predominant site for active corrosion. Such sites where corrosion predominantly occurs in the in-situ corrosion experiments are, of course, the predominant sites for corrosion in the present corrosion-oxidation process resulting in being covered with relatively stable products, because of high pH value in (10 % NaCl + 10 % NaOH) solution. The resultant surface layer formed by the corrosion-oxidation method is considered to prevent the immediate evolution of hydrogen bubbles in the early stage of in-situ corrosion observation and to prolong the time for occurring the filiform corrosion, as shown in (b) and (d), respectively. (a)

(b)

(c)

(d)

200 μm

Figure 3. In situ laser microscopic observations on corrosion behaviors in AZ31 magnesium alloy in 1% NaCl solution. (a) and (c): un-treated specimen. (b) and (d): corrosion-oxidation treated specimen. (a) and (b) show reactions on the surfaces on the un-treated and treated specimens, respectively, after being immersed into the solution. Dark spherical contrasts are due to hydrogen bubbles. (c) shows filiform corrosion occurred on the untreated specimen after about 900 s. (d) shows filiform corrosion on the treated specimen after about 18 ks.

Surface layer in the corrosion-oxidation treated specimen was observed by TEM on a foil prepared by the scratch method as shown in Fig. 4. The selected area diffraction pattern inserted in (b) is composed of diffraction spots and rings, which were analyzed as due to metallic-magnesium (h. c. p.: a0 = 0.321, c0 = 0.521 nm) and magnesium oxide (MgO: f. c. c., a0 = 0.421 nm), respectively. The dark field image in (b) was taken with 200MgO ring. Fine grains of MgO with about 10 to 20 nm in diameter were observed to be formed on the surface. Thickness of the MgO layer seems to be about 500 nm in (b), but the foil was prepared by scratching which may result in wedge-shaped foil, so the apparent thickness of the surface layer would be larger than the proper one. Preparation of foils for TEM observations was difficult to be accomplished mainly because water should not be used. Figure 5 shows an example of microstructures in the foil prepared by a microtome. The dark field image (b) was taken with 200MgO ring, in which broken line indicates the interface between MgO and metallic-Mg. Magnesium oxide film was about 100 nm in thickness, and any defects such as cracks or voids were not observed at the interface.

584 When the specimen was directly oxidized without corrosion, similar microstructures of the surface layer was formed as shown in Fig. 6, in which the specimen was heated at 673 K for 3.6 ks in air. However, the time for occurring the filiform corrosion on the directly oxidized specimen was about 1.8 ks [2]. It is not obvious in Fig. 6 but it is considered that oxidation proceeds into interior region of the specimen with time in the case of direct oxidation, while such an excess oxidation is suppressed by the prior corrosion treatment. (a)

_

(b)

01 ,12Mg

200MgO

200MgO

500 nm Figure 4. Microstructures in the corrosion-oxidation treated specimen: bright field image (a) and dark field image taken with 200MgO ring (b).

(a)

(b)

100 nm Figure 5. TEM observation of the corrosion-oxidation treated specimen of Az31 alloy prepared by means of a microtome. Bright field image (a) and dark fiiedl image take with 200MgO ring. (a)

(b)

(c)

(d)

200MgO

0111Mg

2021Mg

220MgO

2112Mg

200 nm

Figure 6. Microstructures in a specimen directly oxidized at 673 K for 3.6 ks, (a). Dark field images shown in (b) and (c) were taken with 0111Mg spot and 200MgO ring, respectively, showing in the diffraction pattern (d).

585 (a)

(b)

(c)

200 nm Figure 7. Dark field images on the corrosion-oxidation treated specimens of 3N-Mg, AZ61 and AZ91E alloys, (a), (b) and (c), respectively. The images were taken with a part of 200MgO ring.

The corrosion-oxidation treatment on the specimens of 3N-Mg, AZ61 and AZ91E magnesium alloys resulted in formation of thin MgO layer, as shown in Fig. 7.

4

Summary

Microstructures formed on the magnesium alloys by the corrosion-oxidation technique, proposed by the present authors, have been investigated by means of TEM, and it has been concluded that the thin, about 100 nm in thickness, magnesium oxide layer improve the corrosion resistance of the specimens. Origins for corrosion, such as grain boundary segregation, second phase particle, etc., would be covered by corrosion products when the specimens are treated by the corrosion process in this technique and they change into magnesium oxide during the heating process

5

References

[1] Tsubakino, H.; Yamamoto, A. Fukumoto, S.; Watanabe, A.; Sugahara, K.; Inoue, H., Mater. Trans., 2003, 44, 504-510. [2] Yamamoto, A.; Tsubakino, H., Mater. Trans., 2003, 44, 511-517. [3] Yamamoto, A.; Tsubakino, H., Mater. Sci. Forum, 2003, 419-422, 903-908. [4] Maker, G. L.; Kruger, J., Inter. Mater. Review, 1993, 38, 138-153. [5] Ghali, E. Mater. Sci. Forum, 2000, 350-351, 261-272. [6] Yamamoto, A.; Tsubakino, H., “Materials Week 2002, Proceedings”, Ed. Werkstoffwoche-Partnerschaft GbR, Werkstoff-Informationsgesellschaft mbH, Frankfurt, 2003, mw 2002_373.

586

Corrosion and Oxidation of Thixomolded Magnesium Alloys Claudia Pieper, Uwe Köster, Helena Alves, Isao Nakatsugawa* University of Dortmund (Germany); *The Japan Steel Works, Ltd., Hiroshima (Japan)

1

Introduction

AM50, AM60 and AZ91 are standard magnesium alloys which can either be produced by common casting methods like die casting or by the newer thixoforming processes [1]. In the thixoforming processes the alloy is heated to a semi-solid state at temperatures between liquidus and solidus and then injected into a mold. Mechanical shearing of the semi-solid alloy generates a thixotropic structure. The thixoforming processes are distinguished from the exact way of casting. Thixomolding is a process similar to plastic injection molding where the semi-solid mold is sheared by the rotation of a screw. Due to the different temperature of the melt the two production methods thixomolding and die casting will lead to significant differences in the macrostructure (casting defects and porosity) as well as microstructure (different phases and their distribution). Due to the lower temperature before casting thixomolded parts are subjected to less shrinkage, leading to less casting defects and porosity and by that improving the precision of the cast product. Due to the solid portion of the thixoformed alloy the microstructure contains equiaxed nodules of large primary α-Mg surrounded by a finer eutectic mixture of α-Mg and phases much comprised of Mg and the alloying elements; the composition of the remaining melt before casting is enriched in the alloying elements thus leading to a larger volume fraction of the formed second phase or even different phases. It is known that differences in the microstructure not only influence e.g. mechanical properties but also corrosion and oxidation [2]. The aim of this paper is to study the influence of thixomolding as compared to die casting on the oxidation and corrosion behavior. In earlier papers by Nakatsugawa et al. [3] and Mathieu et al. [4] a higher corrosion resistance was found for thixoformed specimen as compared to die cast ones.

2

Experimental

Thixomolded (th) AM60 and (th) AZ91 as well as die cast (dc) AZ91 were prepared by “The Japan Steel Works, Ltd.” and die cast (dc) AM50 was bought from “Arge Metallguss”. The samples were ground and polished up to 3 μm. For investigating the influence of the casting scale only the edges were polished. Before and after oxidation or corrosion the samples were weighed out with a microbalance (Mettler Toledo MT5). For corrosion experiments the samples were immersed for 3 days into a stirred 5% NaCl solution (150 ml/cm2 specimen surface). The 8 h oxidation experiments at 430 °C were done by means of thermogravimetriy (TGA) (Netzsch STA 409) under synthetic air while the experiments at 250, 300 and 400 °C were done in an oven in environmental air.

587 The as-cast specimen and the corroded or oxidized specimen were investigated by means of X-ray diffraction using Cu-Kα radiation, optical microscopy, SEM (Hitachi S-4500 with integrated EDX) and TEM (Philips CM200 operating at 200 kV). For optical microscopy and SEM the samples were ground, polished up to 3 μm and then etched in 1% HNO3. The samples for TEM were either dimpled and ion beam milled (GATAN Duomill 600) in a liquid nitrogen cooled stage or – for cross-sectional microscopy in order to investigate the formed surface layers - thin slices were cut with a diamond knife in an ultramicrotom (Reichert Ultracut S) similar to the preparation of polymers [5] but - different to the procedure for polymers - without water in order to avoid deterioration by corrosion.

3

Results and Discussion

3.1

Microstructure

The microstructure of the as-cast material was investigated in detail. Similar influences of the casting technique used were observed in the AM as well as AZ91 alloys (and will be shown in particular for the AZ91 samples). The most significant differences between die cast and thixomolded alloys are the coarse Mg crystals not dissolved during the thixomolding process (figure 1) and the almost uniform grain size of the newly crystallized Mg grains in thixomolded alloys; die cast alloys show a broad grain size distribution. With decreasing dimension of the mold, leading to thinner parts, the grain size of die cast alloys decreases.

Figure 1. Microstructure (optical microscopy) of die cast (a) and thixomolded (b) AZ91 in the interior of a 2 mm thick plate

Figure 2. Microstructure (SEM) of die cast (a) and thixomolded (b) AZ91 in the interior of a 2 mm thick plate

588 In th AZ91 the cell structure is finer. The small Mg crystals are fully surrounded by a second phase just forming a continous net (in contrast to dc AZ91, figure 2) as already reported in literature [3, 6]. The volume fraction of the second phase is higher as in the die cast alloy. Electron diffraction of this grain boundary phase can be indexed mainly by the β-phase (Al12Mg17) (figure 3a). To some extent quasicrystals can be found (figure 3b) with varying quasilattice constants.

Figure 3. Electron diffraction of the grain boundary phase in thixomolded AZ91 (0.5 mm plate): (a) β-phase; (b) quasicrystalline phase

A gradient in the distribution of large Mg grains towards the center of the plate was found in the cross-sections of the 2 mm thick samples (figure 8b); at the surface there are no large primary Mg grains. In thixomolded pieces, in particular in the thin structures of AZ91 shown in figure 4, typical flow patterns can be found exhibiting significant differences in their microstructure. The eutectic cell size (Mg grains surrounded by Al12Mg17) changes as indicated. There is an irregular distribution of the mentioned large Mg crystals correlated to the flow pattern. Areas with higher flow rate and shear contain less large Mg grains.

Figure 4. Microstructure of three areas in a 0.5 mm thick thixomolded AZ91 plate. Areas were chosen between the large Mg crystals

3.2 Oxidation Oxidation was studied by TGA in synthetic air at 430 °C; the weight gain after 8 h oxidation is shown in figure 5. The oxidation rate of thixomolded alloys is much higher than that of the die cast ones. From equilibrium phase diagrams one can expect that above 400 °C all the β-phase is dissolved into α-Mg, in the die cast as well as in the thixomolded alloys. Dissolution of Al12Mg17 during annealing for 16 h at 410 °C in thixomolded AZ91 was mentioned by Luder et al. [7], but no details on the kinetics were given. This reaction interferes with the ongoing oxidation process. The dissolution of the β-phase, however, is expected to lead to some new porosity along the grain boundaries. This might be the rea-

589 son for the faster oxidation of thixomolded alloys which exhibit a larger amount of β-phase. The higher oxidation resistance of th AM60 as compared to th AZ91 (with the higher Al content) might be related to its smaller volume fraction of the β-phase. At very few singular spots catastrophic oxidation was found to originate (see figure 6) leading to decrepitation of large areas of the specimen. 0.7

24.7

25

0.6

dc AM50 0.5 2 Δm [g/m ]

2 Δm [g/m ]

20 15

7.9

10

0.4 0.3

th AM60

0.2

5

0.2

0.9

1.3

dg AM50

th AM60 Gate

0.1 0

0 dg AZ91

th AZ91 Gate

th AZ91 Overflow

0

50

100

150

200

250

300

350

time [d]

Figure 5. Oxidation of Mg alloys (th: 2 mm p lates, dc: 3 mm plates): (a) weight gain after 8 h oxidation at 430 °C in synthetic air; (b) oxidation kinetics of AM at 300 °C

Figure 6. Catastrophic oxidation of thixomolded AM60 (47 d at 400 °C)

Oxidation studies at 300 °C reveal a faster oxidation of the die cast alloys (see figure 5b for AM). The main difference results from the period of the first half year. Corresponding microstructural changes for AZ91 at 300 °C are shown in figure 7: only partial dissolution of Al12Mg17 and spheroidization at later stages. We assume that in the beginning of the oxidation the higher volume fraction of the β-phase in the thixomolded alloys results in the formation of a more protective scale (with higher Al content).

Figure 7. Microstructural changes during annealing at 300 °C: (a) dc AZ91, 10 d; (b) th AZ91, 10 d; (c) th AZ91, 621 d

At 250 °C no significant differences in oxidation resistance between thixomolded and die cast AZ91 was observed; AZ91 undergoes a slightly slower oxidation than AM.

590 3.3

Corrosion

Corrosion was investigated by means of immersion tests in 5% NaCl at room temperature for different times. Results are shown in figure 8a. 5.6

6.0

4.0

2.8

2.7 2.1 1.6

2.0 1.0

2nd series

3.0

1st series

corrosion rate [mm/y]

5.0

25°C 72h 5% NaCl

0.0 dc AZ91

th AZ91

dc AM50

th AM60

Figure 8. Corrosion of thixomolded and die cast magnesium alloys: (a) immersion test; (b) cross-section of corroded th AZ91 (upper side of the sample with scale, lower side without scale)

The corrosion rate of die cast alloys is significantly higher than of the corresponding thixomolded alloys. Compared to AZ91, the AM alloys exhibit always a much higher corrosion rate. It also can be seen that the (overall) corrosion rate of the specimen with original casting skin is higher than that of the polished specimen. Corrosion rates were observed to decrease with increasing immersion time. The composition of the solution changes during corrosion due to the dissolved corrosion products, thus increasing the pH-value until it reaches a constant value as mentioned already by Lunder et al. [7]. The decreasing corrosion rate might be caused by this increase of the pH-value; changing the solution every 24 h during the three day experiment leads to a constant corrosion rate. The lower corrosion rates of thixomolded specimen can fully be explained by the microstructure. The finer eutectic cell structure of thixomolded alloys with a higher volume fraction of β-phase leads to a continous net of Al12Mg17. In accordance to literature [3, 8] the Mg grains corrode first leading - in case of the fine microstructure of thixomolded alloys - to a new surface consisting of the more noble Al12Mg17 [8]. In die cast alloys the β-phase will fall out when undermined by corroded Mg grains. The scatter in the corrosion rates between different series of specimen are probably due to differences in the macrostructure (e.g. pores), slightly different content of contaminating elements and - in the case of the thixomolded specimen - the influence of the local distribution of the large primary Mg crystals. Whereas the polished surface of AZ91 as well as AM alloys undergo homogenous corrosion independent on the casting method, the natural surface exhibits severe localized corrosion with an overall rate even higher (figure 8b). There is strong evidence that some areas do not undergo any corrosion. The reason for this promising behaviour is still unknown. The fast corrosion in some areas is not related to the local microstructure and might be due to the formation of local cells caused by surface contaminations.

591

4

Conclusion

Oxidation and corrosion behavior of thixomolded and die cast AZ91 and AM alloys was investigated in detail and discussed in regard to the microstructures involved. Thixomolding leads to a microstructure consisting of large primary Mg-crystals embedded in a fine eutectic structure exhibiting a continuous net of the Al12Mg17 phase. During die casting only separated particles of this phase are formed at the grain boundaries. Thixomolding was observed to improve the oxidation resistance, but only at temperatures at or below 300 °C; at 400 °C or higher temperatures oxidation is even accelerated, probably due to the formation of voids during the dissolution of the continuous β-phase at the grain boundaries. The corrosion rate for thixomolded Mg alloys is lower than for the die cast ones. The fine eutectic structure with the continuous β-phase works as a barrier against further corrosion attack.

5

Acknowledgement

The authors like to thank the Deutsche Forschungsgemeinschaft (German Research Foundation) for the financial support under the project Ko 668/25. The authors are indebted to Ms. Monika Meuris for the transmission electron microscopy.

6

References

[1] R. Carnahan, R. Decker, D. Ghosh, C. VanShilt, P. Frederick, N. Bardley, Magnesium Alloys and Their Appl. 1992, 69-76 [2] H. Alves, U. Köster, D. Eliezer; Magnesium 97 (Eds.: E. Aghion, D. Eliezer), Magnesium Research Institute, Beer-Sheva, Israel, 1997, 202-207 [3] I. Nakatsugawa, F. Yamada, H. Takayasu, T. Tsukeda, K. Saito, Proc. of the International Symposium on Environmental Degradation of Materials and Corrosion Control in Metals (Eds.: M. Elboujdaini, E. Ghali), Canadian Institute of Mining, Metallurgy and Petroleum, Montreal, Canada, 1999, 113-123 [4] S. Mathieu, C. Rapin, J. Hazan, P. Steinmetz, Corrosion Science 2002, 44, 2737-2756 [5] L. C. Sawyer, D. T. Grubb, Polymer Microscopy, Chapman and Hall, London, 1987, p. 91 [6] E. Evangeista, S. Spiragelli, P. Cavaliere, Key Engineering Materials 2000, 188, 139-148 [7] O. Lunder, J.E. Lein, T.Kr. Aune, K. Nisancioglu, Corrosion 1989, 45, 741-748 [8] G. Song, A. Atrens, M. Dargusch, Corrosion Science 1999, 41, 249-273

592

Plasma Electrochemical Synthesis of Protective Coatings on WE54 Magnesium Alloy A. Berkania, P. Skeldona, G.E. Thompsona, K. Shimizub, H. Habazakic, H. Karimzadehd, K. Stevense, C. Johne a

b c d e

Corrosion and Protection Centre, University of Manchester Institute of Science and Technology, P.O. Box 88, Manchester M60 1QD, UK University Chemical Laboratory, Keio University, 4-1-1 Hiyoshi, Yokohama 223, Japan Graduate Engineering School, Hokkaido University, N13 W8, Kita-ku, Sapporo 060-8628, Japan Magnesium Elektron, P.O. Box 23, Swinton, Manchester M27 8DD, UK Poeton Industries Ltd., Eastern Avenue, Gloucester GL4 3DN, UK

1

Introduction

Plasma Electrolytic Oxidation (PEO), an emerging new technology for aluminium and titanium, has long been used for the surface treatment of magnesium alloys. Commercial anodizing methods such as the HAE [1], DOW 17 [2] and Cr-22 [3], which have been developed in the 1950s, use sparking to form hard ceramic-type layers on magnesium alloys which provide the metal with an excellent corrosion resistance, especially when the coating is sealed. The recently developed MAGOXID [4], TAGNITE [5] and KERONITE [6] processes employ pulsed currents and involve sparking. Anodizing above the sparking voltage results in visible coloured sparks moving over the electrode, macroscopically covering its surface. The resultant oxide film is usually uniform in appearance and has a complex chemical composition which incorporates elements from the anodizing bath. Because of this, PEO, or spark anodizing, is seen as a promising technique with the potential of tailoring the coatings for specific applications. In the present work, the role of the bath formulation in the DC galvanostatic anodizing of WE54-T6 alloy (composition: Y 5%w, rare earths 3%w, Zr 0.7%w) under sparking was studied in model aqueous single and multi-component electrolytes at 25 °C. The electrolytes were made up by adding to 3 M ammonia one or more compounds at a concentration of 0.05 M each. The compounds were selected from the group: tri-ammonium or tri-sodium orthophosphate, potassium silicate, sodium aluminate, sodium tungstate, potassium fluoride. The range of current -2 densities and forming voltages used were 1–5 Adm and 300–500 V respectively. Sequential anodizing experiments were carried out to provide insight into the mechanism of growth. The morphology and composition of the final 10–30 μm thick films have been characterized in the scanning electron microscope with energy dispersive X-ray analysis. Elemental depth profiles and maps were obtained from electron probe microanalysis. Thin films grown in the initial pre-breakdown stage have been studied using Rutherford backscattering spectrometry.

593

2

Influence of the anion in single-component electrolytes

In preliminary investigations, sodium or ammonium orthophosphates, phosphoric acid, sodium tungstate, potassium or sodium silicate, sodium aluminate, potassium or ammonium fluoride and hydrofluoric acid have been successfully used in solution with ammonium hydroxide or ethanolamine to spark anodize high purity magnesium and commercial magnesium alloys. In order to investigate the influence of the anodizing bath composition, electrolytes containing one (single-component baths) or more (multi-component baths) compounds selected from the above list have been used. Visual observation showed that the anion present in the electrolyte influenced the anodizing behaviour of the WE54 alloy in many respects. First, the type and characteristics of the sparking depended on the electrolyte. For example, in the case of the phosphate compounds, above ≈240 V, uniform fine sparks were observed on the anode surface, the density of which decreased with increasing voltage. The maximum voltage reached was higher than 500 V. Electrolytes containing silicate on the other hand (single- or multi-component electrolytes), lead to the production of an initial glow of the surface with the appearance of the sparks occurring at a later stage. In all electrolytes, variation of the voltage with time at constant current showed two or more different stages. For example, in the case of phosphate, a rapid initial increase of the voltage was followed by a linear increase and a final stage of very slow increase. The voltage corresponding to this plateau-like region of the voltage-time curve was characteristic of the electrolyte used (figure 1). The appearance and uniformity of the coating was also strongly dependent on the electrolyte as was the final thickness reached. Coatings produced in silicate baths were generally smoother than those formed in phosphate baths, but they were also thinner. Less uniform coatings were produced in tungstate baths compared with phosphate, silicate or aluminate baths. The coatings were grey in color, but aluminate baths resulted in greywhite coatings. EDX analyses revealed that the coatings contained Mg, Y, rare earths, O and elements incorporated from the anions used in the bath, 2 namely P for phosphate, Si for silicate, Al for aluminate, F for fluoride and W for tungstate. Elements from the cationic species in the electrolyte were not detected. Except for those grown in baths containing a fluoride compound, EPMA depth profiling showed that the coatings had a uniform composition across their thickness. As an example, the distribution of Mg, Y and P in a coating formed in a phosphate electrolyte is shown in figure 2. A fall in the X-ray signal can be seen which is due to the internal porosity of the layer. Y from the substrate is present throughout the layer.

594

3

Two-component electrolytes

The voltage-time behaviour of the WE54 alloy anodized in 2-component electrolytes containing 0.05 M tri-ammonium orthophosphate is shown in figure 3. The addition of 0.05 M of a given compound to a standard phosphate electrolyte resulted in a change in the position of the ‘plateau’ on the voltage-time curve and in the rate of voltage increase. Thus, tungstate lowers the maximum voltage and the rate of voltage increase, while silicate and aluminate lead to a higher rate of voltage increase. The curves of figures 1 and 3 also show that the influence of the electrolyte composition is relatively insignificant in the pre-breakdown stage of the process below ≈200 V. Again, the composition of the layers incorporated alloying elements from the substrate and two elements from the anions present in the electrolyte. These elements were uniformly incorporated across the layer as figure 4 shows in the case of a phosphatetungstate bath. Porosity of the inner region of the 30 μm layer is revealed here, with local falls in the X-ray signal.

4

Fluoride-containing electrolytes

Fluoride-containing baths formed a different class of electrolytes. In these electrolytes, indeed F was present in the layer, but mainly at the interface with the substrate (figure 5). This suggests that in fluoride baths F-reacts with Mg to form a thin layer of MgF2 at the interface with the metal. This thin magnesium fluoride layer is shown in the EPMA map of figure 6 relating to a coating grown in a fluoride electrolyte containing phosphate and silicate.

5

Sequential anodizing

Experiments were conducted where samples were anodized to 400 V in a first singlecomponent anodizing bath and then reanodized to a maximum voltage of ≈500 V in a second single-component bath containing a different anion. EPMA depth profiling showed that in this case, elements were incorporated in the coating from the anions in both electrolytes. The element from the second bath was always predominantly present in the central part of the coating. Thus, the initial layer formed in the first bath has been displa-

595

Fig. 1. Influence of the anion on the voltagetime behaviour of WE54. Current density: -2 -2 1 Adm (3 Adm for tungstate).

Fig. 2. EPMA elemental line profiles across the final coating formed in a phosphate electrolyte. -2 Current density: 1 Adm . Forming voltage: 490 V.

Fig. 3. Voltage-time behaviour of WE54 in 2-component electrolytes containing phosphate. Comparison is made with a single-component phosphate electrolyte (1). Current density: 3 Adm-2.

Fig. 4. EPMA elemental line profiles of a coating formed in a phosphate-tungstate bath. Current density: 3 Adm-2. Forming voltage: 410 V.

Fig. 5. EPMA elemental line profiles showing the distributions of P and F in a coating formed in a phosphate-fluoride bath. Current density: -2 1 Adm . Forming voltage: 480 V.

Fig. 6. EPMA elemental map showing the distribution of F in a coating formed in a 3-component bath -2 containing fluoride. Current density: 1 Adm . Forming voltage: 460 V.

596

Fig. 7. Switching from a silicate to an aluminate bath. EPMA depth profiles. Current density: 3 Adm-2. Forming voltage in first bath: 400 V (80 s). Forming voltage in second bath: 500 V (280 s).

Fig. 8. Switching from a phosphate bath to a phosphate-fluoride bath. EPMA depth profiles. Current density: 3 Adm-2. Forming voltage in first bath: 400 V (123 s). Forming voltage in second bath: 490 V (478 s).

Fig. 9. Switching from a silicate-fluoride bath to a phosphate bath. EPMA depth profiles. Current density: 3 Adm-2. Forming voltage in first bath: 400 V (70 s). Forming voltage in second bath: 490 V (356 s).

ced during spark anodizing in the second bath through formation of new material at the interface with the substrate. The decrease of the concentration of the element incorporated from the first bath towards the outermost part of the layer also indicates formation of new material at the coating-electrolyte interface during reanodizing. Figure 7 illustrates this phenomenon in the case of a transition from a silicate bath to an aluminate bath. A question then arises on whether fluorine would be incorporated again as a layer of MgF2 after reanodizing in a fluoride bath a sample that has been previously anodized in a fluoridefree electrolyte. Indeed, as figure 8 shows, reanodizing in a phosphate-fluoride bath a sample

597 that has been anodized to 400 V in a phosphate electrolyte resulted in the formation of a fluorine-rich thin layer at the interface with the metal. The behaviour of the fluoride ion provided a convenient way to assess the mechanism of growth of the coating and was used to confirm the previous results. When a sample was anodized in a silicate-fluoride bath, reanodizing in a phosphate electrolyte resulted in a magnesium fluoride layer being located in the central part of the coating (figure 9). Some fluorine, however, was retained adjacent to the alloy. Thus, during spark anodizing, new material forms at the metalcoating interface as well as at the coating-electrolyte interface.

6

Correlation between the bath composition and the coating composition

RBS analysis has been carried out on thin films formed in the pre-breakdown stage of anodizing in different electrolytes. The results showed that the level of the elements incorporated from the anionic species in the bath increased with anodizing time and reached a stable value in the first minutes of anodizing. The final element to magnesium atomic ratios are summarized in table 1. These clearly show that the composition of the final coating is determined by the composition of the bath. For example, increasing the concentration of phosphate in a phosphate-silicate electrolyte increases proportionally the amount of phosphorus in the coating and vice versa. When equimolar 2-component baths were used, the element to Mg atomic ratios in the coating were the same for the elements originating from the two anionic species. Table 1. Correlation between the electrolyte composition and the final coating composition.

Anion(s)

(Molarity)

Final coating Voltage (V)

Phosphate (0.05) Silicate (0.05) Phosphate (0.05) Phosphate (0.05) Phosphate (0.1) Phosphate (0.01)

7

490 470 Tungstate (0.05) 409 Silicate (0.05) 480 Silicate (0.01) 430 Silicate (0.1) 460

Element/Mg at. ratio Time (s)

P

1342 1215 543 902 – 1114

0.719

Si

W

0.525 0.558 0.577 0.759 0.082

0.556 0.588 0.030 0.862

Conclusion

(1) Spark anodizing of WE54-T6 alloy at constant current density in alkaline electrolytes containing phosphate, silicate, tungstate and aluminate as a single anion or a mixture of anions results in the formation of coatings 10–30 μm thick that are uniform in composition throughout their thickness. The coatings incorporate elements from the substrate and the electrolyte. (2) Fluorine, in fluoride-containing baths, incoporates as a thin layer of magnesium fluoride at the metal-coating interface.

598 (3) Sequential anodizing experiments confirmed that during spark anodizing new coating material forms at both interfaces. (4) There is a correlation between the relative concentrations of the anionic species in the electrolyte and the final chemical composition of the coating which makes it possible to formulate the anodizing bath in order to achieve specific compositions of the layer.

References [1] H. A. Evangelides, Electrolytic coating of magnesium., Patent US2723952 (1955). [2] Anonymous, Bath for and method of producing a corrosion resistant coating upon light metals., Patent GB762195 (1956). [3] W. McNeill, Corrosion-resistant coatings., Patent US2778789 (1957). [4] E. L. Schmeling, B. Roschenbleck, M. H. Weidemann, Method for finishing magnesium and magnesium alloys., Patent EP333049 (1989). [5] D. E. Bartak, B. E. Lemieux, E. R. Woolsey, Hard anodic coating for magnesium and magnesium alloys prepared by two-step process., Patent US5470664 (1995). [6] A. S. Shatrov, Method for producing hard protection coatings on articles made of aluminum alloys., Patent WO9931303 (1999).

599

Galvanic Corrosion of Magnesium Alloy AZ91 with Coated Elements E. Boese1, J. Goellner1, A. Heyn1, St. Eisenberg2 1 2

Otto-von-Guericke Universität, Magdeburg, Germany Volkswagen AG, Wolfsburg, Germany

1

Introduction

In using magnesium alloys, special attention has to be paid to protection against galvanic corrosion which arises at various fastener locations required by design in automobile components [1, 2]. For example there were investigations to start serial production of the magnesium hatchback for the 3 l car [3]. General observations on galvanic corrosion are provided in ASTM G71 [4], DIN 50919 [5] and by Schwenk [6]. In order to assess the corrosion susceptibility of the less noble element in a material combination, it is necessary to know the characteristics of the partial polarization curves of both materials in the aggressive medium, not just the size of the potential difference between the two connected metals. The extent to which anodic or cathodic partial reactions can take place unhindered, or are restrained (e.g., by coatings), is decisive for the extent of galvanic attack. A coupling current, which increases the rate of corrosion of the anode, flows in a galvanic couple [4, 5]. However this current is not directly equivalent to the corrosion rate, since it does not take into account the natural corrosion occurring at the anode. The natural corrosion rate of the anode is that part which exceeds the rate arising from the coupling current and is caused by the cathodic partial reaction occurring at the anode. The coupling current is a complex parameter: it is dependent upon the geometrical arrangement, the size of the electrode surfaces in contact with the medium, the free corrosion potentials and the polarization resistances of the two metals, as well as the electrolytic resistance of the medium. Simplified consideration of galvanic corrosion processes is only possible under the assumption of homogeneous distribution of current. The characteristic parameter for electrochemical corrosion processes is the coupling-current density which can be transformed into a loss in weight via Faraday’s law.

2

Experimental Procedures

Determination of the coupling current density was the focal point of this research into galvanic corrosion. In addition measurements of polarization resistance and electrochemical noise were made to characterize the properties of the coatings before and after contact with AZ91. A sodium chloride solution (1N NaCl; 25 ºC, pH 7.0) was used as electrolyte.

600 2.1

Galvanic Corrosion according to DIN 50919

2.1.1 Specimen Materials Galvanic Corrosion of the alloy AZ91, containing 9% Al and 1% Zn, was investigated in contact with different fasteners. These were covered with four coating systems to minimize galvanic corrosion: 1. Duplex coating system: zinc/nickel coating + cathodic immersion paint 2. Duplex coating system: zinc coating + cathodic immersion paint 3. Thick galvanized zinc coating (ca. 20 μm) with a silicate sealant 4. Galvanized zinc + yellow chromate + organic coating 2.1.2 Experimental Method An experimental equipment with homogenous distribution of the galvanic current was used. The two electrodes to be examined were placed in the electrolyte in a constant distance from each other. The short circuit to form a galvanic element was produced by an external electrical connection, within which the galvanic-coupling current could be directly measured. The measured potentials were converted to the standard hydrogen electrode (SHE) scale. After each test, the corrosion products were removed from the anode with chromic acid. The electrodes were dried and weighed before and after the experiment. The difference of the rate of weight-loss per unit area (vtot), calculated from the weight loss of AZ91 coupled to a coated part, and the rate of weight-loss per unit area (vi), calculated from the coupling current density of the galvanic element, was compared with natural corrosion of AZ91 without contact under the same conditions. 2.2

Measurements of Impedance and Polarization Resistance

Impedance and polarization resistance measurements were made on various coated fasteners (coats 1, 3, 4) in a three electrode cell. The coated fasteners were characterized after exposure at the free corrosion potential without contact to AZ91 and these values were compared with the coating resistances following galvanic coupling to AZ91. The polarization resistance was measured by polarizing to +/- 50 mV from the free corrosion potential with a polarization rate of 1 mV/second and determining the gradient of the anodic and cathodic branches to the curve. Impedance measurements were carried out in the frequency range from 100 kHz to 10 mHz (resolution 10 data points per decade) with an amplitude of 10 mV. Software was used to generate the Nyquist plot and the necessary curve fit for determination of R1, R2 and C, which were simulating the electrolytic resistance (R1), the coating resistance (R2) and the double-layer capacitance (C). 2.3

Electrochemical Noise Measurements

Electrochemical noise measurements were carried out at a selected couple of AZ91/coat 1 to demonstrate the effect of defects in the coat after a short time. These measurements were done with a standard three electrodes setup by use of a zero-resistance ammeter and voltmeter. The upper frequency was limited to 1 Hz and the readings were taken with 2 Hz during the whole time by a 16 bit AD converter. For a better resolution of the electrochemical noise the DC part was separated from the measuring device and the AC part was amplified by a factor of 100. The resulting effective frequency range of the electrochemical noise was 0,01-1 Hz.

601

3

Results

3.1

Galvanic Corrosion according DIN 50919

3.1.1 Protecting effects of various coatings

galvanic coupling current density i

rate of weight-loss per unit area v [g/m²d]

20 18

c

[A/m²]

During 168 h in contact with AZ91 the coating types 3 and 4 showed no significant changes in the galvanic coupling current density (Figure 1). But it was observed that adhesion to the zinc undercoat had become lost. The coatings could be stripped off. coat 1

16

coat 3

14

coat 4

12 10 8 6 4

125 v tot [g/m²d] v i [g/m²d] 100

75

50

25

2 0

0 0

24

48

72

96

120

144

168

time t [h]

a) time dependence of current densities

AZ91 natural dissolution

coat 1

coat 3

coat 4

b) rate of weight-loss per unit area

Figure 1. AZ91 in contact with hold-down devices having various protective coatings after 168 h; a) Time dependence of galvanic coupling current densities, b) Material consumption rate per unit area calculated from galvanic coupling current densities (vi) and weight loss of the anode(vtot)

However, contact of AZ91 with components protected by coating 1 showed a steep increase in the coupling current density and rate of weight loss per unit area after about 24 h, which indicated increasing corrosion. The difference between vtot and vi is much larger than the value of the natural corrosion rate of the magnesium alloy without contact. Thus, the natural dissolution (intermetallic corrosion) of magnesium alloys is also increased by strong galvanic corrosion. According to [7], this intermetallic corrosion of magnesium alloys results from electrochemical corrosion and then proceeds in parallel without measurable external current. During galvanic corrosion, the formation of hydrogen bubbles is observed both on the anode and on the cathode. After 168 h of contact with the component covered with coating 1, the anode (AZ91) was very strongly attacked and, in fact, had nearly dissolved away; the coating itself showed blisters and spalling. 3.1.2 Influence of the Coating Quality The quality of coatings is of decisive importance when in contact with magnesium alloys. This will be demonstrated by the behaviour of various components protected by a coating consisting of cathodic immersion paint of different quality and layer structure (coatings 1 and 2). Independent of the geometry and size of the component, good protection against galvanic corrosion is initially provided for AZ91 in contact with components of 2 coating type 1 (Figure 2a). The galvanic coupling density lies between 0.2 and 0.4 A/m at 2 times till 12 hours, whereas it is one decade lower at 0.03 A/m for contact of AZ91 with components of coating 2. However, after about 12 till 24 hours in contact with coating 1, the coupling current density strongly increases and galvanic corrosion of AZ91 continues apace. In contrast, the coupling current density remains low over the whole time period of

100

rate of weight-loss per unit area v [g/m²d]

galvanic coupling current density ic [A/m²]

602 coat 2 screw coat 1 threaded bushing coat 1hold-down device coat 1plug

10

1

0.1

125 vtot [g/m²d] 100

vi [g/m²d]

75

50

25

0

0.01

AZ91 coat 2 coat 1 coat 1 coat 1 natural screw bushing plug hold-down dissolution 0

24

48

72

96

120

144

168

time t [h]

a) time dependence of current densities

b) rate of weight-loss per unit area (168 h)

Figure 2. Galvanic corrosion of AZ91 with components coated with cathodic immersion paint of different quality; a) time dependence of current densities, b) material consumption per unit area calculated from galvanic coupling current densities (vi) and weight loss of the anode (vtot)

the experiment when in contact with parts covered by coating 2, showing that this coating system provides complete protection against galvanic corrosion. The differences in protective quality between the coating systems 1 and 2 were confirmed by the determination of rates of weight loss per unit area. No galvanic corrosion of the magnesium alloy took place when in contact with components covered by coating 2 (Figure 2b). The rate of weight loss per unit area corresponded to that for natural corrosion of AZ91. In contrast, the protective effect of the coating 1 got lost. 3.2

Polarization Resistance measurements

Figure 3 shows the polarization resistance values of components with different coatings, both after exposure at the free corrosion potential and when in contact with AZ91.

1E+6 1E+5 1E+4 1E+3 1E+2

coat 1 1,5 h

coat 3 168 h

coat 4

within contact to AZ91

polarization resistance [ohm x cm²]

polarization resistance [ohm x cm²]

without contact to AZ91

1E+7

1E+7 1E+6 1E+5 1E+4 1E+3 1E+2

coat 1 1,5 h

coat 3

coat 4

168 h

Figure 3. Polarization resistance of hold-down devices with various protective coatings after different test times; a) after exposure at the free corrosion potential, b) after contact with AZ91

603 It is apparent that the components protected by coating 1 show a resistance after exposure at the free corrosion potential (Figure 3a) which is three decades higher than that for the fasteners protected by coatings 3 and 4. No major effect results from different exposure times. After contact of the fastener to AZ91, the polarization resistances of the components covered with coating 1 are reduced by three decades, irrespective of the time of exposure (Figure 3b). In contrast, the resistance of coatings 3 and 4 remains in the same order of magnitude as during exposure at the free corrosion potential. 3.3

Measurements of Electrochemical Noise

current [mA]

The results which were obtained by determination of the coupling current density will now be confirmed and examined more precisely by measurements of electrochemical noise. In order to determine the exact point in time at which the protective system really fails, the time course of the current, the electrochemical current noise and the potential of the coupled element were determined simultaneously for contact of AZ91 and fasteners with coating 1 (Figure 4a). The coupling current increases very slowly and a noticeable change is only apparent after about 24 hours. However, the first characteristic signals for galvanic current noise are registered already after about 5 hours. Thereafter, the noise signals increase continuously. The potential of the galvanic couple is in the region of the free corrosion potential of the magnesium alloy (ESHE = -1310 mV). Without contact to AZ91, the corrosion potential of the fastener with coating 1 is much more positive (ESHE = -435 mV). The result of this change in potential of the cathode while in contact with the magnesium alloy is over-polarization, which favours the formation of hydrogen blisters at the cathode. As well as causing rapid dissolution of the magnesium anode, contact with coatings of low quality also leads to the simultaneous destruction of the protective coating system. Figure 4b shows the initial destruction of the top coat after 48 hours contact of the part with coating 1 and AZ91. Defects in the protective coating can be detected very much earlier with the help of electrochemical noise measurements than it is possible by conventional measurements of the coupling current alone. 5 4 3 2 1

potential U SHE[mV]

current noise [µA]

0 2.5 2 1.5 1 0.5 0 -0.5 -1 -1275 -1300 -1325 -1350 -1375 -1400 -1425 -1450

0

12

24

36

48 time /h

0

12

24

36

48 time /h

0

12

24

36

48 time /h

a) Transients of electrochemical noise measurements

b) Defects on the top coat

Figure 4. a) Time dependence of current, current noise and potential of the galvanic element AZ91 in contact with a cathodic coated joining element with lower quality (coat 1); b) Destruction of the top coat of the cathodic coated joining element with lower quality (coat 1) by bubble formation and sub-surface corrosion after 48 h galvanic contact

604

4

Conclusions

- At high contact corrosion rates the difference between vtot and vi has been much larger than the natural corrosion of AZ91. This means that the quota of the natural dissolution also rises with increasing galvanic corrosion. - In contact with AZ91, the potential of the coated components is displaced into the region of the free corrosion potential of AZ91, i.e. to strongly negative potentials. Overpolarization occurs, which favours the formation of hydrogen blisters at the cathode. As well as causing rapid dissolution of the magnesium anode, contact with coatings of low quality also leads to the simultaneous destruction of the protective coating system. - The coating 2 on steel components prevented galvanic corrosion when in contact with AZ91 over the whole period of testing. This top had the best protective properties. Steel components coated with top 1 caused a drastic increase of galvanic corrosion connected to a drop in the coating resistance. Coatings 3 and 4 prevented galvanic corrosion quite well during one week. However, at the end of the experiment it was observed that adhesion to the zinc undercoat had become lost. The protective coating could be stripped off. - Results from impedance and polarization resistance measurements support the results from galvanic corrosion testing. - Defects in the protective layer can be detected very much earlier and more accurately with the help of electrochemical noise measurements than by conventional measurements of the coupling current alone.

5

References

[1] SAE 2001 world congress: "Light metal applications for the automotive industry aluminum and magnesium", SAE SP-1622, Detroit, Michigan, USA, 05.-08.03.2001. [2] H. Schreckenberger, "Korrosion und Korrosionsschutz von Magnesium-Werkstoffen für den Automobilbau - Problematik der Kontaktkorrosion", Düsseldorf, FortschrittBerichte VDI, Reihe 5 613, 2001. [3] H. Schreckenberger, M. Papke, and S. Eisenberg, "The magnesium hatchback of the 3-liter car: processing and corrosion protection", presented at SAE 2000 World congress, Detroit, Michigan, 2000. [4] "Conducting and evaluating galvanic corrosion tests in electrolysis", ASTM G81, 1981. [5] "Korrosion der Metalle - Korrosionsuntersuchungen der Kontaktkorrosion in Elektrolytlösungen", DIN 50919, 1984. [6] W. Schwenk, "Probleme der Kontaktkorrosion - Aufgaben der Materialprüfung für die Anwendung", Metalloberfläche, vol. 35, pp. 158-163, 1981. [7] H. Haferkamp, "Magnesiumkorrosion - Prozesse, Schutz von Anode und Kathode", in Magnesium - Eigenschaften, Anwendungen, Potentiale, K. U. Kainer, Ed. Weinheim: DGM, Wiley-VCH Verlag GmbH, 2000, pp. 244-259.

605

Corrosion Protection by Conductive Polymers for Magnesium Alloys A. Yfantis1*, I. Paloumpa1, D. Schmeißer1, D. Yfantis2 1 2

BTU Cottbus, Department of Applied Physics and Sensors, Cottbus, Germany NTU Athens, Department of Chemical Engineering, Athens, Greece

1

Introduction

The advanced corrosion prevention properties of conducting polymers are known and established, but their application is limited due to the lack of application methods with industrial potential. Our group has developed novel methods for application of conducting polymers on aluminum substrates [1] and has attempted to modify these techniques properly in order to apply them on magnesium surfaces, too. Our preliminary studies revealed the importance of the Mg oxide, the limitations of the in situ polypyrrole film formation [1] and the significant potential of the technique based on the film formation from a polyacrylic-polypyrrole emulsion (“PACPY” coatings). The polyacrylic emulsion had also been studied for the preparation of corrosion resistant films on tinplate [2]. In this work, we present the experimental data related to the effect of the polypyrrole concentration and the film thickness on the corrosion inhibition properties of these novel polymer coatings. The relative corrosion resistance properties of the films formed from the matrix solutions, where polypyrrole was in situ polymerized and in cases where polypyrrole is dispersed in a form of powder (pigment) in the polyaclylic polymer, are further examined. The evaluation of the experimental data provides significant evidence regarding the mechanism of corrosion inhibition, as well as the optimization of the synthesis and application conditions.

2

Experimental

Immersion tests were performed in a corrosive environment of 1% w/v NaCl at neutral pH (6.9) for a time period of 9 days. The samples had been prepared by using 7 different emulsions of a polyacrylic copolymer and polypyrrole (Ppy). Magnesium samples (Mg AZ31) 2 have been used with a total exposed surface area of 400.5 mm . Three samples of each coating have been tested as well as three samples without any coating (reference samples). At first, in all cases the coatings were formed after double immersion in the polyacrylicpolypyrrole matrix solution for 10 seconds. A time interval of 5 min was allowed to pass between the first and the second immersion. Finally the samples were dried at 80 C for 15 min (Table 1 - samples 4A to 4D1).

606 The thickness of the dry films were further examined by using the solution providing the coating with the best performance. The new samples had different film thickness obtained with intermediate drying and re-coating several times. In this case the samples were dried for 5 min at 80 C before every re-coating process (Table 2 - samples 5A to 5D). Table 1. Mg AZ31 samples coated with different PACPY coatings. Sample code

4A

4B

Solution No 23% v/v (aqueous, treatment polyacrylic emulsion always two immersions)

Coating weight

–

0.17 2 mg/cm

4B1

4B2

4C

4C1

4D

23% v/v polyacrylic emulsion and 0.010 g/ml Ppy (in situ)

23% v/v polyacrylic emulsion and 0.020 g/ml Ppy (in situ)

32% v/v polyacrylic emulsion

32% v/v 100% v/v polyacrylic polyemulsion acrylic and emulsion 0.010 g/ml Ppy (in situ)

100% v/v polyacrylic emulsion and 0.030 g/ml Ppy (powder, PTS doped)

0.41 2 mg/cm

1.06 2 mg/cm

0.29 2 mg/cm

0.39 2 mg/cm

1.48 2 mg/cm

0.88 2 mg/cm

4D1

The corrosion rate was monitored every second day with the determination of the concentration of Mg in the solution by using AAS (Atomic Absorption Spectrophotometry). The corrosion morphology was examined by using an optical videomicroscope with a magnification up to 1000x. The density of the polymer 4C1 was determined experimentally to be 0.85 g/cm3 when the density of the polyacrylic polymer itself is 0.70 g/cm3. The conductivity of the 4C1 polymer was determined by using the 4 probe technique on thick films (20 - 30 μm) formed on glass surfaces. The conductivity was 0.01 S/cm and less than 10-6 S/cm for the acrylic polymer itself.

3

Results

In Figure 1. the average corrosion rate for samples of the acrylic polymer itself (different film thickness) as well as the blank sample (no coating) are shown. It is evident that increasing the film thickness of the acrylic polymer itself by a factor of 5 (see Table 1) can only reduce the corrosion rate by 15 % (samples 4B & 4D). Table 2. Mg AZ31 samples coated with the best performing PACPY coating of different thickness.

Sample code 5A Solution (aqueous)

Coating weight

5B

No treatment 32% v/v polyacrylic emulsion and 0.010 g/ml Ppy (in situ); one immersion –

2

0.20 mg/cm

5C

5D

32% v/v aqueous polyacrylic emulsion and 0.010 g/ml Ppy (in situ);

32% v/v aqueous polyacrylic emulsion and 0.010 g/ml Ppy (in situ);

three immersions

four immersions

2

0.53 mg/cm

0.66 mg/cm2

607 The average corrosion rate of the overall test period for all the samples is shown in Figure 2. The sample 4C1 had the best performance with a corrosion rate reduction -when compared to the uncoated one- up to 70%. The surface density of the coating was relatively low (0.39 mg/cm2), when all coatings had been in the area of 0.2 -1.4 mg/cm2. Samples with different concentrations of polypyrrole (4B, 4B1, 4B2, 4C and 4C1) provided a significant reduction of the corrosion rate upon addition of polypyrrole and increasing its content. The reduction was 27% and 47% in the case of samples with increasing polypyrrole content (4B1 and 4B2) when compared with the sample coated only with the acrylic copolymer (4B). Even greater was the reduction of the corrosion rate of the samples with increased copolymer concentration (C 1) where the reduction was 60% when compared with the sample with no polypyrrole addition (C). The film thickness dependence of the average corrosion rate for the best performing PACPY coating is shown in Fig 3. The corrosion rate reduction is only 9% when the film thickness increased by a factor of 3.

2

corrosion rate ( g/mm h)

0,25

0,2

0,15

0,1

0,05

0 0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1

coating weight (mg/cm2)

Figure 1. Pure polyacrylic film thickness dependence of the corrosion rate.

The samples with and without the addition of Ppy powder (4D and 4D1) indicated that polypyrrole addition as a pigment has a very limited effect on the corrosion rate. Samples D and D1 had a marginal difference of 9.5 % although their surface density has a difference of a factor of 2 (Fig.4).

4

Discussion

Corrosion Rate [µg/mm²h]

The resulting film thickness of the polymer during the preparation is actually related to both:

0.16 0.14 0.12 0.1 0.08 0.06 0.04 0.02 0

4A

4B

4C

4D

4B1 4B2

0%

23.50%

4C1

32%

0 0.01 0.02 100%

Ppy [g/ml]

polyacrylic [v/v]

Figure 2. Corrosion rate dependence of Mg samples coated with polyacrylic – polypyrrole film (PACPY) of different thickness and composition (all samples according to Table 1).

608 the viscosity of the matrix solution in which the metallic surface is immersed, as well as the drying process. The coating weight is directly correlated to the film thickness, when polymers with the same specific weight are compared. For example, samples 4B, 4C and 4D have been formed from the same acrylic polymer in different dilution ratios (viscosity). From the Table 1 it can be seen that they have significant differences in the film thickness. It can be safely assumed that sample 4D has a polyacrylic coating as much as 5 times thicker than sample 4B. However, the reduction of the corrosion rate is only 15% indicating that the polyacrylic polymer itself has no corrosion inhibition properties (Fig.1). Consequently the significant reduction of the corrosion rate upon polypyrrole addition can be exclusively attributed to the action of polypyrrole. One of the main aspects to be addressed is the mechanism of corrosion inhibition. Regarding the corrosion inhibition offered from conducting polymers there are four mechanisms widely discussed in the literature [3]: • The conducting polymer may act as a barrier coating due to the repulsion of the negative ions from the negatively charged dopants surrounding the polymer chain • The conducting polymer may function as an anode and provide cathodic protection to the metal surface corrosion rate (μg/mm2 h)

0,25

0,2

0,15

0,1

0,05

0 0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

2

coating weight (mg/cm )

Figure 3. Dependence of the corrosion rate on the PACPY film thickness (best performing PACPY coating).

• The conducting polymer may act as a reservoir of ions, turning the corrosive anions into dopants or releasing the original dopants which subsequently act as inhibiting ions • The conducting polymer could stabilize a passive oxide layer The last mechanism has been extensively used in order to explain the corrosion resistance of polymeric conducting layers (mainly polyanniline) [4], especially in cases where the corrosion inhibition action was evident even in the non conductive state of the polymer. However, there is no further justification or evidence of the mechanism of formation of these passive films. In our case the film has a very low conductivity and it is very thin (up to 5 μm) to provide any barrier properties (normally expected for film thickness greater than 60 μm). Furthermore, in Fig. 3 one can see that increasing the film thickness of the best performing PACPY coating had finally a detrimental effect on the corrosion rate. In a previous study [1] we have established that a strong chemisortpion process of pyrrole monomers or oligomers may take place on Al, Ti or Zr oxides. The chemisorption process has evident from a 0.9 eV shift of the Nitrogen N 1s peak of polypyrrole at the

609 XPS spectra. In this case we have established that pyrrole could stabilize the structure of the oxides and provide advanced corrosion properties. Consequently, a similar mechanism can be suggested in the case of the PACPY coating. This mechanism is further verified considering that retaining the native oxide on the Mg surfaces had a positive result on the corrosion inhibition properties of the applied film. Furthermore as shown in Fig.4, the addition of polypyrrole in the form of a pigment has a very limited effect on the corrosion rate, indicating that the protective action of polypyrrole is minimized when the polymers are not in the physicochemical form to interact with the metal interface. Our data clearly demonstrate that the reduction of the corrosion rate can be as much as 70% and it is exclusively attributed to polypyrrole. The ratio of pyrrole and polyacrylic polymer in the matrix solution is an important factor for the resulting corrosion inhibition properties of the film. Raising the polyacrylic concentration from 23,5% (4B1) to 32% (4C1) and keeping polypyrrole content and other synthesis conditions constant, the reduction of the corrosion rate when compared to the uncoated sample was boosted from 30% to 70%. It can be assumed, that there is an interaction between the acrylic polymer and the polypyrrole, when prepared in situ, which also affects on the final structure of the dry film. Since it is known that in the conventional preparation of polypyrrole the concentration of the dopant can significantly affect on the properties of the resulting conductive polymer, it can be assumed that the acrylic copolymer may act as a dopant.

4C

4D

corrosion rate 2 [μg/mm h]

0.14 0.12 0.1

4C1

0.08

4D1

0 0.01

0.06 0.04 0.02 0

0.02

Ppy [g/ml]

0.03 32%

100%

polyacrylic [v/v] Figure 4. Comparison of coatings with polypyrrole polymerized in situ and added as a pigment (All samples according to Table 1).

Consequently, increasing the polypyrrole ratio itself is equally important as optimizing the ratio and most probably other synthesis conditions which will allow the acrylic copolymer to become dopant of polypyrrole.

610

5

Conclusions

The in situ polymerization of pyrrole in an aqueous polyacrylic emulsion according to the developed method can give rise to an aqueous solution, where a thin anticorrosive coating can be formed on a Mg alloy surface after simple immersion and drying. The optimized synthesis conditions had as a result the achievement of the reduction of the corrosion rate in respect to the uncoated specimens up to 70% . The corrosion rate of the coated samples showed only a weak film thickness dependence, indicating that the bulk properties of the coating are not of great importance. Furthermore the addition of polypyrrole to the polyacrylic matrix in the form of a powder (pigment) provide thick films with reduced corrosion resistant properties. We argue that the corrosion inhibition properties of the films should be attributed to the formation of a thin passive layer in between the polymeric film and the surface metal oxides. This layer is formed from polypyrrole polymers or oligomers strongly chemisorbed on the metal oxides. The structure is further stabilised from polyacrylic polymer chains interacting as dopants of polypyrrole. The properties of this layer are affected from the composition and synthesis conditions of the polyacrylic – polypyrrole emulsion. The use of non toxic chemicals, no etching, no rinsing and therefore no effluents make the method environmental friendly. Additionally, the advanced corrosion inhibition properties of the films, as well as the simple application process on the metal surface provide a significant potential for the industrial application of the method.

6

References

[1] A.D. Yfantis, D. Schmeißer, D.K. Yfantis, L. Kompotiatis, J. ATB Metallurgie 2000, XXXX, Nr. 3-4; J. ATB Metallurgie 2001, XXXXI, Nr 1-2, 527. [2] D. Yfantis, A. Yfantis, B. Tzalas, D. Schmeißer, Corrosion ( NACE ) 2000, 56, No.7, 700. [3] V.T. Truong, P.K. Lai, B.T. Moore, R.F. Muscat, M.S. Russo, Synthetic Metals 2000, 110, 7. [4] N. Ahmad, A. G. Mac Diarmid, Synthetic Metals 1996, 78, 103.

611

Plasma Hydrogenation of Mg-based Alloy Films Under Highflux, low Energy ion Irradiation at Elevated Temperatures L. L. Pranevicius1,2, D. Milcius2, G. Thomas3 1

Vytautas Magnus University, 8 Vileikos St., LT-3035 Kaunas Lithuanian Energy Institute, 3 Breslaujos St., LT-3035 Kaunas 3 Sandia National Laboratories, 537 Spirit Ridge Court, Reno, NV 89511 2

1

Introduction

In recent years, has been an increasing interest in the use of magnesium and magnesium based alloys as hydrogen-storage materials, because they can store more hydrogen by weight than most of the other currently know metal hydrides (MgH2 is equivalent to 7.6 wt.% hydrogen). It becomes clear that in the new generation of advanced materials a proper engineering of the alloy composition, surface properties, microstructure, grain size etc. is needed to design and control the adsorption-desorption properties of hydrogen [1,2]. Nanocrystalline hydrides are of a great interest [3,4]. Thin film physical vapor deposition technologies allow to apply non-traditional and new nanotechnology methods for designing materials according to the requirement of particular application. Plasma assisted deposition is widely used to synthesize complex chemical compounds [5]. Hydrogen atoms can be incorporated in a controlled manner into growing film employing simultaneous vapor deposition of metal atoms (Mg and Al) and hydrogen ion irradiation technique. In this case, not only the interaction effects modify new phases synthesis kinetics and structure of surfaces. Energetic ions penetrate into the sub-surface monolayers of growing film and saturate it by the incident hydrogen atoms. These materials are usually not thermodynamically stable. They contain many grain boundaries, defects, disorder or strain. The properties of these materials cannot be fully described by the basic thermodynamic functions, so that behavior of hydrogen can differ significantly from theoretical expectations. In this situation, we have planned to apply plasma activated physical vapor deposition nanotechnology methods for synthesis of thin magnesium alanate film material and to study the correlation between changes of physical structure and chemical features.

2

Experimental technique

There are many modifications of physical vapor deposition (PVD) processes in preparing thin film and coating materials. In the present work plasma-enhanced magnetron-sputtered deposition (PMD) process, described in [6,7] was used. PMD combines conventional magnetron sputtering with independently generated plasma that provides for independent control of the sputter-deposition and plasma ion irradiation processes.

612 In the vacuum chamber (0.5 m diameter and 1.2 m long) two cylindrical magnetrons (7.0 cm diameter) symmetrically surrounded the substrate holder that hosted samples 1 x 1 cm2 in size made of stainless steel. All samples were ultrasonically cleaned using acetone then methanole. The gap distance between magnetrons and sample was 5 cm. A dc discharge power supply was used to produce plasma that was independent for the plasma produced by the magnetrons. This independently generated plasma enveloped the sample holder, which was biased negatively, to perform the ion-sputter cleaning before the deposition and implantation of hydrogen during the deposition. The cathode of magnetron was made of Mg and Al with a nominal purity better than 99.99 at.%. The sample temperature during the deposition process was found to be an important parameter that affected the synthesis kinetics. Therefore, to ensure the accuracy of the temperature measurement, a thermocouple was spot-welded to the back-side of sample holder. After having been cleaned, the samples were installed in the vacuum chamber. A typical argon sputter-cleaning process was used to remove the surface contaminants and the residual oxide. The magnetrons were not used in this process. A bias voltage of 100 V was applied to the sample holder. At the end of the sputter–cleaning process, after 15 min, the co-deposition of Mg and Al atoms process from two independent magnetrons proceeded without any interruption. During deposition, 500 W power on Mg magnetron and 600 W power on Al magnetron were used. The partial pressure of hydrogen was controlled in such way that concentration of hydrogen in growing film would be sufficient or more which was needed to synthesize Mg(AlH4)2. The supply of hydrogen incorporated into the growing film was performed on the basis of the simulation results described in literature [8]. When a typical thickness of approximately 3-5 μm was obtained (in approx.10 min), the Mg magnetron was switched out and the thin about 20 nm thick Al layer was deposited on the top of the film. After exposure of samples in air, thin 2-3 nm thick barrier layer of natural oxide Al2O3 was formed on the surface which protected chemical compounds from the direct contact with air and moisture. Then the samples were removed from the chamber and subjected to the structural and properties analysis. The structural properties have been studied by X-ray diffraction with the 2Θ angle in the range 20-70º using CuKα radiation in steps of 0.05º. The grain sizes were estimated by Debye-Scherrer analysis of the XRD spectrum. The scanning electron microscopy (JOEL JSM-6300) was used for the studies of surface morphology and topography on all stages of fabrication of thin films.

3

Results

3.1

Plasma hydrogenation of Mg films

Fig. 1 shows X-ray diffraction patterns of the as-deposited Mg film (a) and after the hydrogenation in H2 plasma for 30 min at 380 K (b). It is seen that the Mg-film is converted into MgH2. The estimate grain size of crystallites is equal to 30 ± 6 nm. Fig. 2 includes experimental dependencies of the MgH2 phase fractional volume in 5 μm thick Mg film as a function of the hydriding time at different temperatures: 350 K – curve 1; 375 K – curve 2; 400 K – curve 3, and 450 K – curve 4. Other plasma treatment parameters: bias voltage – 100 V, pressure – 10 Pa, ion current – 80 mA. It is seen that the

613 efficiency of hydriding is low at temperatures below 370 K and sharply increases at temperatures above 400 K.

Figure 1. The XRD patterns of Mg films: (a) – as-deposited; (b) – after plasma hydriding for 30 min at 380 K.

Figure 2. Dependencies of the MgH2 hydride volume fraction in Mg film on the hydriding duration for different temperatures: 350 K – curve 1; 375 K – curve 2; 400 K – curve 3 and 450 K – curve 4.

Figure 3. The SEM surface views of Mg films: (a) – as-deposited; (b) – after hydrogenation at 370 K for 30 min.

The threshold hydrogenation temperature is an important parameter in understanding of the hydriding mechanism. The SEM surface analysis shows (Fig. 3) that the as-deposited Mg film has homogeneous surface structure (a), which becomes non-homogeneous with well developed grain boundaries and channels between them after plasma hydriding for 30 min at 420 K.

3.2 Plasma hydrogenation of MgAl films Recent experimental results [9] on the synthesis, structure and thermal properties of magnesium alanate, Mg(AlH4)2, which contains 9.3 wt.% of hydrogen, exhibit promising features as a hydrogen storage material. The investigated material has been obtained using chemical reaction between sodium alanate and magnesium chloride. In this paper the physical vapor deposition technologies are used to synthesize magnesium alanate at low temperature with controllable microstructure and composition.

614 In the PMD, the argon and hydrogen ions of energy controlled up to 500 eV bombard a growing film on the substrate, transfer the ion momentum to the adsorbed Mg and Al atoms, enhance the mobility of atoms on the substrate and new phases synthesis kinetics. The end product of the reaction has been registered by X-ray diffractometry. Fig. 4 includes X-ray diffraction patterns of the films that have been synthesized at different temperatures: 300 K – Fig. 4a, 400 K – Fig. 4b, and 475 K – Fig. 4c. The identification of phases has been performed using Crystallographica Search-Match program.

Figure 4. XRD patterns of MgAl films hydrogenated at 300 K – (a); 400 K – (b) and 475 K – (c) for 30 min.

The mean size of crystallites of Mg(AlH4)2 in the as-deposited for an ion current to the substrate equal to 40 mA film is about 30-50 nm. The size of crystallites does not depend notably on temperature in the temperature range 225-500 K, however, it depends on the intensity of ion irradiation during film growth. The mean size of crystallites decreases up to 20-30 nm with the increase of ion current directed to the growing film up to 120 mA.

Figure 5. The SEM surface views of MgAl films deposited on stainless steel substrates and hydrided during deposition: (a) – 430 K and (b) – 510 K.

615 SEM surface analysis of magnesium alanates is related to some problems. Alanates have low electrical conductivity. The thin 20 nm thick Al film, which was deposited on the top of magnesium alanate film, solved this problem. The presence of the 2-3 nm thick layer of natural oxide Al2O3 did not distort the SEM image. Additionally, it was noticed that SEM images change during e-beam exposure may be due to decomposition of hydride phases under electron irradiation. In order to minimize this effect, beam currents were kept as low as practical while maintaining adequate signal-to-noise ration in the images. Fig. 5 includes SEM surface views of films synthesized at different temperatures: a – 430 K and b – 510 K. The SEM analysis results show that the surface topography manifests about the processes in the bulk. The hydrogenation at elevated temperatures results in the surface topography manifesting the presence of bubbles in the bulk of the film (Fig 5a), which grow with the increase of temperature. At temperatures above 500 K the surface topography of synthesized films reveals the presence of cracks and liftings as is seen in Fig 5b.

4

Discussions

Mg(AlH4)2 contains large amount of hydrogen. However the release of hydrogen does not occur in a single step reaction. The decomposition of magnesium alanate proceeds as follows: Mg(AlH 4 ) 2 → MgH 2 + 2AlH 3

(1)

2AlH3 → 2Al + 3H 2

(2)

The first reaction starts at a temperature between 375 K and 400 K [10]. The decomposition of AlH3 takes place at temperatures above 400 K. Finally the decomposition of MgH2 occurs at a much higher temperatures around 500-520 K [11]. For hydrogen storage, only the second reaction needs to be considered. Reversibility of this reaction is a critical factor for the usability of magnesium alanates in hydrogen storage devices. Reversibility means the ability to form Mg(AlH4)2 directly from MgH2 and Al, under hydrogen plasma treatment. In the present work the magnesium and aluminum in elemental form were used and magnesium alanate was formed by the following reaction Mg + 2Al + 4H 2 → Mg(AlH4 ) 2

(3)

The experimental results prove that MgH2 and Mg(AlH4)2 are synthesized during deposition of Mg and Mg+Al atoms in hydrogen plasma. The XDR and SEM surface analysis results are in agreement with the decomposition reactions of synthesized hydrides. The hydrogen released in MgH2 films is effused out, however, the hydrogen released in Mg(AlH4)2 films is clustered into bubbles. A compact amorphous Al2O3 layer typically 2-5 nm thick limits the hydrogen effusion and acts as permeation barrier for hydrogen.

5

Conclusions

The MgH2 and Mg(AlH4)2 films were successfully synthesized using vapor deposition of Mg and Al atoms in hydrogen plasma under negative 100 V bias voltage. The synthesized films demonstrated nanocrystalline structure. The decomposition of MgH2 occurs at temperatures around 500-520 K.

616 The decomposition of Mg(AlH4)2 proceeds in two stages: (i) at a temperature between 370 K and 400 K magnesium alanate decomposes into MgH2 and AlH3, and (ii) at temperatures above 400 K hydrogen is released and clustered into 10-20 μm diameter bubbles. The effusion of hydrogen out of the film is hindranced by 2-5 nm thick Al2O3 barrier layer.

Acknowledgements The work was supported by the Sandia National Laboratories (USA). Authors express special gratitude for Mr. A. Vasys (USA) for the encouragement and help in an administration of the research, and wish to thank Prof. Liudvikas Pranevicius for helpful discussions.

References [1] A. J. Maeland in Recent Advances in Hydride Chemistry (Ed.: M. Peruzzine), Elsevier, 2001, 531 - 557. [2] G. Sandrock, K. Gross, G. Thomas, J. Alloys and Comp. 2002, 339, 299 - 308. [3] K. Higuchi, H. Kajioka, K. Toiyama, H. Fujii, Y. Kikuchi, J. Alloys and Comp. 1999, 484, 293 - 295. [4] K. Higuchi, K. Yamamoto, H. Kajioka, K. Toiyama, M. Honda, S. Orimo, H. Fuji, J. Alloys and Comp. 2002, 330 - 332, 526 - 531. [5] L. Pranevicius, Coating Technology: Ion Beam Deposition, Satas&Associates, Warwick, Rhode Island, 1993, p. 453. [6] J. N. Matossian, R. Wei, J. Vajo, G. Hunt, M. Gardos, G. Chambers, L. Soucy, D. Oliver, L. Jay, C. M. Taylor, G. Alderson, R. Komanduri, and A. Perry, Surf. Coat. Technol. 1998, 108 - 109, 496 - 509. [7] S. V. Fortuna, Y. P. Sharkeev, A. P. Perry, J. N. Matossian, and A. Shuleopov, Thin Solid Films 2000, 377 - 378, 512 - 526. [8] L. Pranevicius, C. Templier, J. Delafond, and S. Muzard, Surf. Coat. Technol. 1995, 72, 51 - 61. [9] M. Fichtner, O. Fuhr, and O. Kircher, J. Alloys and Comp. 2003 (in press). [10] E. C. Ashby, R. D. Schwartz, B. D. James, Inorg. Chem. 1970, 9(2), 325 - 332. [11] L. Zaluski, A. Zaluska, J. O. Strom-Olsen, J. Alloys and Comp. 1997, 253, 70 - 75.

617

The Corrosion Performance of Sealed KERONITE Coatings on AZ91D Magnesium Alloy S. Hutchinsb, S. Shrestha a, A. Sturgeon a, P. Shashkov b and A. Shatrov b a b

TWI Ltd, Granta Park, Great Abington, Cambridge CB1 6AL, UK Keronite Ltd, Granta Park, Great Abington, Cambridge CB1 6GX, UK

1

Introduction

Magnesium is of growing interest in automotive1, electronics2, aerospace3, textile, printing and other industries, due to its favourable strength to weight ratio. Thin walls, high casting rates and long die life are typical advantages of cast magnesium. A popular alloy is AZ91D (9%Al, 1%Zn), with good mechanical and physical properties and excellent castability4-6, but its poor corrosion resistance limits its use. Corrosion resistance of AZ91D can be improved by electrolytic coating processes like the KERONITE™ process that involves creation of a plasma discharge around a component immersed in a proprietary, environmentally-friendly, low concentration alkaline electrolyte. A modulated voltage is applied to the component, creating sparks at its surface. This causes oxidation of the surface as well as elemental co-deposition from the electrolyte, creating a hard ceramic layer. Coating growth is <4μm/minute and working temperature 20-50 °C.

2

Experimental Approach

2.1

Materials and Test Specimens

The nominal composition of the AZ91D alloy considered in this work is bal. Mg, 9Al, 1Zn, 0.2Mn, max. 0.001Fe, max. 0.001Ni and max. 0.015Cu. in both die cast and thixocast forms. Keronite coatings of 25-30μm thickness were made on 40x40mm specimens. The coatings on die cast magnesium were sealed with sol-gel solution A, while those on thixocast magnesium were sealed with sol-gel solution B. The coating process is commercially sensitive and not discussed in this report. The test specimens and coating types are listed in Table 1. Table 1. Test specimen identification

Magnesium alloy type Diecast Thixocast

Test specimen Magnesium Keronite coating Keronite coating Magnesium Keronite coating Keronite coating

Post coating surface Label condition Uncoated Mg-a As coated KC-a Sealed with sol-gel type A KCSG-a Uncoated Mg-b As coated KC-b Sealed with sol-gel type B KCSG-b

618 2.2

Coating Characterisation

Coated specimens were encapsulated in Araldite, sectioned and then polished to 1 μm diamond finish. The obtained cross sections were examined using a scanning electron microscope (SEM). Cystalline composition of the coating was determined by X-ray diffraction. 2.3

Salt Spray Exposure Test

The salt spray exposure was carried out as described in the ASTM standard B117-97. The surface quality was given a rating number in accordance with ASTM D1654-92.

3

Results

3.1

Coating Characterisation

Figure 1 shows SEM images of the coatings on the two alloys. Some porosity and a few micro-cracks were observed in both coatings. The top surface of the coatings appeared to be rough with several loose particles standing on the surface. Figure 2 shows as-coated surfaces of the Keronite. The surfaces display agglomerated particles formed during the micro-arc oxidation process. The Keronite surface on the die cast specimen (KC-a) in Figure 2a appeared to contain some porosity (dark areas) between the agglomerated particles. The X-ray diffraction confirmed MgAl2O4 as the major constituent with minor amounts of SiO2 and SiP2. Mounting compound

a)

Mounting compound

b)

Coating Coating

AZ91D die cast

AZ91D thixocast

Figure 1. SEM sections of the Keronite coating on: a) die cast; and b).thixocast specimens.

a) Figure 2. SEM images of the Keronite coating on: a) die cast; and b) thixocast specimens.

b)

619 3.2

Salt Spray Exposure

Ratings according to ASTM D1654-92 procedure B (Figure 3), show poor corrosion resistance of the uncoated magnesium surfaces (Mg-a and Mg-b). Initially, thixocast magnesium (Mg-b) displayed better resistance than die cast (Mg-a), as shown by its gradual decay in the rating. However, the surfaces of both uncoated specimens deteriorated significantly after just 24 hours. Figure 4 shows the endurance of all test surfaces to the rating of 9. All the coated and sealed specimens displayed 336-hour endurance. The unsealed Keronite coating on thixocast alloy (KC-b) displayed a 9 rating in excess of 1000-hours and performed better than sealed Keronite on die cast substrate (KCSG-a), which had a 750-hour endurance. Sealed Keronite on thixocast substrate (KCSG-b) endured 2000 hours with no sign of corrosion attack. 12 10

Rating

8

Mg-a KC-a

6

KCSGa Mg-b

4

336 hours

KC-b

2

KCSGb

0 0

200

400

600

800

1000

1200

1400

1600

1800

2000

Time, hours

Figure 3. Rating of the test surfaces exposed to salt spray environment.

Salt spray endurance, hours

2000

1500

1000

500

0 Mg-a

Mg-b

KC-a

KC-b

KCSG-a

KCSG-b

Figure 4. Salt spray endurance to a rating of 9 according to ASTM D1654-92 procedure B

620

a)

Mg-a

d)

1cm

b)

KC-a

1cm

e)

1cm

KC-b

1cm

f)

c)

KCSG-a

Mg-b

1cm

KCSG-b

1cm

Figure 5. Optical images of the coatings after 1000-hour exposure to salt spray environment

621 Surface appearances after a 1000-hour salt spray exposure are shown in Figure 5. The thixocast magnesium (Mg-b) displayed significantly less corrosion. This was also reflected in the improved corrosion resistance provided by the Keronite coating on the thixocast magnesium compared to die cast magnesium. The sealed Keronite coating on die cast magnesium displayed some improvement in the corrosion resistance compared to the unsealed coating. This is reflected by the rating in Figure 3 and the photographs (Figure 5b and c). The sealed Keronite coating on thixocast magnesium showed no evidence of corrosion attack after 1000 hours of exposure. Figure 5 shows a few stains in the form of parallel lines which appear to be corrosion product flowing from the edge over the exposed surface. The surface of the sealed coating on the thixocast specimen displayed no sign of corrosion attack after exposure for 2000 hours as shown in Figure 6, appearing similar to the one observed after 1000-hours in Figure 5f.

KCSG-b

1cm

Figure 6. Optical image of coating KCSG-b after 2000-hour salt spray exposure.

4

Discussion

This work suggested that the coating formed on the thixocast magnesium alloy mainly comprised spinel (MgAl2O4). The SEM images of the coatings suggest that the coating microstructure and the level of coating porosity differ for the two types of alloy. A more porous structure is formed on the die cast alloy. It has been reported that the amount and form of casting porosity depends upon the casting method employed, the design of the casting, and thermal and feeding parameters7 and this may have an effect on the characteristics of the Keronite coating. The presence of a ceramic layer on magnesium alloy would be expected to improve its resistance to corrosion attack. Keronite coating on the die cast alloy gave 336 hour endurance in the salt spray environment to a rating of 9. The improved corrosion resistance of the Keronite coating on the die cast alloy can be seen in Figure 5b. The Keronite coating alone on the thixocast alloy displayed further improvement and gave a 1000-hour salt spray endurance at a rating of 9. An improved corrosion resistance of the Keronite coating on the thixocast alloy can be expected due to the denser coating microstructure.

622 The effect of the post-sealing with sol-gel clearly enhanced further the corrosion resistance of the Keronite coatings on both alloys. By sealing, it was possible to achieve salt spray endurance up to 720 hours at rating 9 for the die cast alloy and more than 2000 hours for the thixocast alloy at a rating of 10. This suggests that the Keronite coating on the sealed thixocast alloy acts as an effective barrier to corrosion attack for long-term exposure.

5

Conclusions

1. Microscopic examination showed that Keronite coating on thixocast magnesium is less porous than on a die cast magnesium. 2. The coating is primarily a ceramic layer of Spinel (MgAl2O4). 3. The Keronite coating can improve the corrosion resistance of both die cast and thixocast alloys in a salt spray environment and can provide a minimum of 336-hour salt spray endurance. 4. Unsealed Keronite on a thixocast alloy can give up to 1000-hour salt spray endurance. 5. Sealed Keronite further improves corrosion resistance in aggressive salt spray environment. 6. By sealing with an appropriate sol-gel the Keronite coating can provide thixocast magnesium alloy salt spray endurance exceeding 2000 hours with no sign of corrosion attack.

6

References

[1] M. Brettmann, B. Reinhold, and H. Schnattinger, The Use of Magnesium in Cars – Aspects of Corrosion, Magnesium Alloys and their Applications, K. U. Kainer ed., Wiley-VCH, Verlag GmbH, 2000, p. 271-273 nd [2] B. Landkof, Magnesium Application in the Electronic Industries, Proc. 2 Israeli Intl Conf. on Magnesium Science & Technology, E. Aghion and D. Eliezer ed., Magnesium Research Institute Ltd, Israel, 2000, p. 50-56. [3] G. B. Evans, Applications of Magnesium in Aerospace, Magnesium Technology, Proc. London Conf., November 1986, Institute of Metals, London, 1987, p. 103-109. [4] R. L. Edgar, Global Overview on Demand and Applications for Magnesium Alloys, Magnesium Alloys and Their Applications, K. U. Kainer ed., Wiley-VCH, Verlag GmbH, 2000, p. 3-8. [5] G. Govender and J. H. Ferreira, Erosion Corrosion of Magnesium AZ91D in the Ascast Condition, Proc. 2nd Israeli Intl. Conf. on Magnesium Science & Technology, E. Aghion and D. Eliezer ed., Magnesium Research Institute Ltd, Israel, 2000, p. 371376. [6] J. I. Skar and D. Albright, Emerging Trends in Corrosion Protection of Magnesium die castings, Magnesium Technology 2002, H. I. Kaplan ed., 2002, p. 255-261. [7] G. Song and A. Atrens, Corrosion Behaviour of Skin Layer and Interior of die cast AZ91D, Magnesium Alloys and their Applications, B. L. Mordike, and K. U. Kainer ed., Werkstoff-Informationsgesellsschaft mbH, Germany, 1998, p. 415-419.

623

Corrosion Properties of PVD-coated Magnesium Die Cast Alloy AZ91D Part II: Contact Corrosion Properties of Differently PVD treated Magnesium Die Cast Alloy AZ91 H. Hoche 1, C.Blawert 2, E. Broszeit 1, C. Berger 1 1

Staatliche Materialprüfungsanstalt und Fachgebiet und Institut für Werkstoffkunde an der Technischen Universität Darmstadt, Darmstadt 2 GKSS Forschungszentrum Geesthacht GmbH, Geesthacht

Abstract The publication is divided into two parts. Part I described the general coating and corrosion properties as well as the electrochemical properties of the following coating systems: • 9 μm CrN hard coating [1] • 3 μm TiN hard coating • 0.5 μm plasma anodisation (PA) layer & 3 μm Al2O3 coating [2] Summarizing the results of part I, the PA+Al2O3 coating system showed outstanding corrosion behaviour in comparison to the CrN and the TiN coatings. In the salt spray test the PA + Al2O3 specimens at least came off as good as a commercially available plasma chemical anodisation with sealed porosity. The current densities vs. potential curves obtained in the electrochemical measurements allow an additional classification of the corrosion behaviour of PVD coated AZ91 using the cathodic brunch of the polarisation curve to calculate the corrosion rate. However further information about the corrosion mechanisms involved can be obtained if the surface and cross section of corroded specimens are studied. TiN and CrN reveal debonding of more or less the whole coating at the interface, while the interface remains intact for the PA + Al2O3 coating. Part II describes the results of a novel immersion technique to determine the time dependent corrosion behaviour of coated magnesium alloys [3]. Furthermore contact corrosion properties of the above mentioned coatings according to DIN 50919 are presented. The specimens were tested in an 3% NaCl solution against the aluminium alloys AlMgSi 0,5 and AlSi9Cu3 and against a roller bearing steel 100 Cr 6. The tested specimens were investigated by means of optical microscopy, SEM and EDX.

1

Results

1.1

Immersion Test at constant pH-value

This is still a quite new procedure for corrosion testing of magnesium alloys originally developed by GM [1], which was now adapted for characterisation of coating systems on magnesium by the authors (GKSS). The coated corrosion specimens and as a reference the

624 substrate material AZ91 were exposed to 5% aqueous NaCl solution at pH 6. Any increase in the pH value due to Mg corrosion at defects in the layer was neutralised by adding 0.01 m HCl according to GM test [3]. By recording the amount of used HCl as a function of time the corrosion rate was calculated from the slope of the curve, giving the corrosion rates over the period of the experiment. More details about the test and the data evaluation will be presented elsewhere [4]. Additional information about the major corrosion mechanisms were obtained by observing the surfaces of the specimens with a stereo microscope while being immersed in the solution. Figure 1 shows the time dependent corrosion rates obtained from the immersion test method and figure 2 shows the photographs of the specimen surfaces in several stages during the test runs. For comparison the corrosion rate of the uncoated and polished AZ91 is included in figure 1. Figures 1 and 2 indicate that only the PA-Al2O3 layer can provide a sufficient corrosion protection for the AZ91 alloy. Even after 50 hours of contact with the NaCl solution only minor defects of the layer can be detected. However the corrosion rate of the PVD-PAAl2O3 is very low during the test run but it is not constant. At the beginning the corrosion rate falls from 2 mm/a to nearly 0 mm/year after 20 h testing time and after 50 h the corrosion rate is in the range of 1,5 mm/year. Even after 50 h no significant corrosion products are forming on the surface although hydrogen bubbles evolve on the surface from time to time. The Al2O3 layer is not conducting and appears to be less noble in comparison to the substrate, thus even defects in the layer do not enhance the corrosion rate and formation of a quasi passive film at coating defects containing hydroxides of aluminium and magnesium obviously occurs. This is supported by the observation of active areas (hydrogen evolution) which turned into passive areas after a while. To clarify this behaviour more detailed investigations and longer testing times are necessary. Both, the CrN and TiN coatings enhance the corrosion rates of the substrates in comparison to the AZ91. However there are distinct differences in the corrosion mechanism. Starting with a corrosion rate of approx. 29 mm/year for the CrN coated surface the rate rapidly decreases and after 15 h the slope flattens so that a stationary corrosion rate of approx. 9 mm/year is reached. In contrast the corrosion rate of the TiN coated specimen increases with testing time from 6 mm/year in the beginning to 15 mm/year after 24 hours. This poor corrosion behaviour can be related to the large potential difference of coating and substrate and to a reasonable electrical conductivity of the coatings. Thus existing defects in the layers will become immediately active sides when the electrolyte reaches the substrate and no passivation is possible. Corrosion starts from small pits forming at the already existing defects and while there is continuous grow of the pits in depth and width for the TiN layer, the CrN layer corrodes very fast. In later case the electrolyte migrates underneath the layer and hydrogen evolution and/or formation of corrosion products is so high that whole chips from the coating were blasted from the surface. After 8 h of testing almost 95% of the layer is removed and the corrosion rate is approaching the corrosion rate for uncoated AZ91. After the fast removal of the CrN layer the corrosion is quite uniform, while for the TiN localised pitting corrosion is the dominant mechanism.

625

corrosion rate in mm/year

30 25 Cr-CrN

20 15

Ti-TiN

10 AZ91 5 PA-Al2O3 0 0

10

20

30

40

50

time in h Figure 1: Corrosion rates obtained in the immersion tests.

1.2

Contact Corrosion Properties

The contact corrosion properties were tested according to DIN 50919 [5] in 3% NaCl solution. The contact materials were extruded aluminium alloy AlSiMg0,5, aluminium die cast alloy AlSi9Cu3 and a roller bearing steel 100Cr6. The spacing between the samples was 50 mm. The contact corrosion measurement duration was 4,5 h. before and after the measurements the free corrosion potentials were measured for a duration of 15 minutes. For data logging a measuring amplifier type Hottinger MGC, controlled by a PC with the CATMAN 2.1 software, was used. A general rule in the constructional sector for avoiding contact corrosion is to avoid large differences in electrochemical potentials between the contact partners. Therefore the AlSiMg0,5 alloy should show better contact corrosion properties than the copper alloyed AlSi9Cu3 and this should again show better contact corrosion properties than the roller bearing steel 100Cr6 [6]. Table 1 shows the free corrosion potentials for the used coating systems and for the contact materials, measured in 3% NaCl solution. The duration of the measurements is 5 h. Table 1 includes the free corrosion potentials after 10 minutes and after 5 h.

626

Figure 2: Sample surfaces during the immersion tests

627 Table 1: Corrosion Potentails in 3% NaCl solution after 10 minutes and 300 minutes.

Time

10 min 300 min (5h)

PVD coated AZ91 specimens

contact materials

PA+Al2O3

Cr-CrN

Ti-TiN

AlMgSi0,5

AlSi9Cu3

100Cr6

-1,3 V -1,29 V

-1,27 V -1,25 V

-1,3 V -1,29 V

-0,64 V -0,52 V

-0,48 V -0,44 V

-0,31 V -0,39 V

The evolution of the corrosion potentials of the PVD coated specimens and the uncoated and polished AZ91 in 3% NaCl solution during 5 h testing time is plotted in Figure 3. As can be seen, the corrosion potentials of the TiN and the CrN coated specimens are nearly constant and smooth and the values are comparable with the uncoated alloy. This implies that there is several coating porosity so that a mixed corrosion potential dominated by the less noble substrate is measured. The corrosion potential of the PAAl2O3 coated sample increases from -1400 mV in the very initial state of the measurement to -1280 mV after 15 minutes. The run of the curve is not as smooth as for the TiN and the CrN coatings with a deviation of approx. 20 mV from the mean value. This may be related to the formation of corrosion products at existing coating defects. Another mechanism could be a steady interplay between corrosion and passivation in microscopic scale. Free Corrosion Potentials in 3% NaCl -1200 -1220

Cr-CrN

-1240

Potential in mVH

-1260

PA-Al2O3

-1280

Ti-TiN

-1300 -1320

AZ91D

-1340 -1360 -1380 -1400 0

2000

4000

6000

8000

10000

12000

14000

16000

18000

Time in s

Figure 3: Evolution of the free corrosion potentials of the PVD coated AZ91 specimens in 3% NaCl for 5 h.

The diagram in Figure 4 contains the integrated contact currents normalized to the couple AZ91D-AlMgSi0,5 according to the presentation in [7]. The value for the couple AZ91D-AlMgSi0,5 is set to 1. Higher values indicate higher integrated corrosion currents and therefore worst contact corrosion properties. For all coating systems it is evident that the higher the differences in free corrosion potential between coated AZ91 and contact material are, the higher are the resulting values for the normalized integrated corrosion currents in figure 4. The superior corrosion

628 properties of the PA-Al2O3 coatings in comparison to Ti-TiN and to Cr-CrN coatings are no more valid under contact corrosion conditions. The diagram in figure 4 indicates the Ti-TiN coatings as best contact corrosion partners. Even the value for the couple TiNAlMgSi0,5 is lower than for the couple AZ91D-AlMgSi0,5 which implies that the Ti-TiN coated AZ91 has better contact corrosion properties than the uncoated alloy itself. The values for the contact corrosion couples with the Cr-CrN coated specimens are in the same range and even better than the PA-Al2O3 coated ones. However as we have to assume defects in the coatings we have a complex galvanic system. Magnesium can dissolve not only due to the Al or steel, but also due to the coating around the defect and also Al and Fe can dissolve due to the more noble TiN and CrN coatings. All this can take place at the same time and can contribute to the measured current. In any case and this includes also the PA-Al2O3 coating no passivation of the magnesium is possible anymore, as all contact partners have a more noble potential than magnesium in NaCl solution. Comparing to the polarisation measurements in part I of the paper, the magnesium is always in the anodic branch of continuous dissolution of magnesium. Nevertheless the TiN and CrN coatings are more noble than the contact partners Al or Fe, thus they can protect as long as they cover the surface. As indicated by the immersion tests only the TiN coating last over a longer period while CrN is blasted from the surface by enhanced magnesium corrosion. The potential of Al2O3 appears to be similar or less noble than AZ91 and therefore no protection can be expected under contact conditions. To evaluate the attack on the magnesium specimens requires further analysis. 5,06

CrN vs. 100Cr6 3,50

CrN vs. AlSi9Cu3 1,17

CrN vs. AlMgSi0,5 TiN vs. 100Cr6

4,73

TiN vs. AlSi9Cu3

2,27

TiN vs. AlMgSi0,5

0,41

PA-Al2O3 vs. 100Cr6

5,78

PA-Al2O3 vs. AlSi9Cu3

3,61

PA-Al2O3 vs. AlMgSi0,5

1,30

PA-Al2O3 vs. AZ91D

0,10

AZ91D vs.AlMgSi0,5

1,00

0

1

2

3

4

5

6

Figure 4: Integrated contact corrosion currents normalized to the couple AZ91D-AlMgSi0,5.

To get an impression of the overall corrosion attack on the PVD coated surfaces, the corrosion products were removed by etching in hot Cr6+ acid (Figure 5). Although the complexity of the system the corrosion attack of the coated samples shows strong dependence onto the ratios of the normalized contact corrosion currents from Figure 4. As expected, the material with the lowest difference in free corrosion potentials, AlMgSi0,5, causes the lowest corrosion attack on the PVD treated AZ91 surfaces and vice versa the 100Cr6 causes the highest damage. The corrosion attack of the Ti-TiN coated surfaces is considerable the lowest. Both the PA-Al2O3 and the Cr-CrN show similar corrosion attack for each contact material. While with the roller bearing steel 100Cr6 the attack onto the Ti-TiN coated surface is only selective, the PA-Al2O3 and the Cr-CrN coatings are crucially affected by corrosion.

629 Cr-CrN

Ti-TiN

100Cr6

AlSi9Cu3

AlMgSi0,5

PVD-PA-Al2O3

Figure 5: Photographs of the PVD coated AZ91 specimens after contact corrosion testing. Diameter of tested area is 14 mm.

2

Conclusion

The following list gives the conclusions obtained from the investigations presented in Part I: Investigation of the Corrosion Properties and the Corrosion Mechanisms of PVD Coated Magnesium Die Cast Alloy AZ91 for Tribological Applications and “Part II: Contact Corrosion Properties of Differently PVD treated Magnesium Die Cast Alloy AZ91”: • Due to the ionic character of the Al2O3 coatings and remarkable back-sputtering effects the maximum coating thickness within an acceptable deposition time is 3 μm. The adhesion of the PA-Al2O3 coatings is acceptable. The best adhesion values showed the Cr-CrN coatings while the adhesion of the Ti-TiN coatings is weak. • Comparing the overall corrosion performance in the salt spray test shows the superiority of the PA-Al2O3 coatings in comparison to the Cr-CrN and the Ti-TiN. Even a surface damage does not lead to a deterioration of the corrosion properties. The plastic deformation underneath the scratch seems to improve the corrosion resistance. A possible reason therefore can be the densification of the intermetallic Mg17Al12 network.

630 • The corrosion performance in the salt spray test of the pre-damaged PA-Al2O3 coated sample is at least as good as for a pre-damaged conventional plasma chemical anodised AZ91 sample. • Current densities vs. potential curves confirm a superior corrosion rate for the PAAl2O3 coated samples. The corrosion rates can be calculated from the cathodic branch using the Tafel equation. The anodic behaviour is comparable for all coating systems and allows no differentiation. Investigation of the samples in the corrosion initiation state discloses crucial differences between the coating systems where the PA-Al2O3 nearly shows no damage and the Cr-CrN and Ti-TiN show significant coating detachment and corrosion in the interface between coating and substrate. • The immersion tests confirm the superiority of the PA-Al2O3 coatings. Defects are always present in the PVD coating thus passivation of magnesium in the defect is a crucial point. The results show that only PA-Al2O3 coatings allow passivation of the magnesium with improved corrosion resistance. All other coatings enhance the corrosion rate in comparison to uncoated AZ91. TiN coatings suffer from severe pitting corrosion while CrN coating are completely removed from the AZ91 surface due to interface failure. Chips of CrN are blasted from the surface by strong hydrogen evolution. • None of the investigated coating types is able to provide sufficient contact corrosion protection. Here the Ti-TiN coatings did the best job while both the PA-Al2O3 and the Cr-CrN coatings were crucially affected by contact corrosion. TiN and CrN fail by more or less the same mechanisms as in the immersion tests, thus the galvanic couple coating/AZ91 seems to be dominant. TiN suffers only from pitting so only a small amount of Mg is exposed to the electrolyte and might be also affected by the more noble contact partner. CrN is not stable and chips are falling of the surface exposing large areas of pure AZ91. The potential of Al2O3 appears similar or less noble than AZ91, so less protection under the present contact conditions can be expected and passivation of defects is not possible anymore. Therefore the TiN appears to be the best choice for contact corrosion. • However looking at the overall excellent corrosion performance of the Al2O3 coating a sealing top coating should be considered if it will be used under contact corrosion conditions. TiN and CrN appears to be no alternative from a corrosion point of view.

Acknowledgement The authors would like to acknowledge the financial support by the DFG (Deutsche Forschungsgemeinschaft) for the studies presented in Part I and Part II of this publication.

631

Literature [1] H. Hoche, H.J. Schroeder, H. Scheerer, E. Broszeit, C. Berger, “Tribological Studies of CrN coated Magnesium AZ91 at Temperatures up to 250 °C”, Advanced Engineering Materials 4, (2002), 42-51 [2] H. Hoche, H. Scheerer, E. Broszeit, C. Berger., „Erschließung von Mg-Legierungen für tribologisch-korrosive Komplexbeanspruchung mittels durch Plasmaanodisation erzeugter Korrosionsschutzschichten“, Mat.-Wiss. u. Werkstofftech., 33, (2002), 395-403 [3] B.L. Tiwari and J. J. Bommarito, “A novel technique to evaluate the corrosion behaviour of magnesium alloys”, Magnesium Technology 002, TMS, (2002), 269-275 [4] C. Blawert, W. Dietzel, K. Pfeiffer and V. Heitmann, “A new method for corrosion testing of coating systems on magnesium alloys”, to be published [5] DIN 50919, „Korrosionsuntersuchungen der Kontaktkorrosion in Elektrolytlösungen“, Beuth Verlag GmbH, Berlin (1984) [6] ASM Specialty Handbook, „Magnesium and Magnesium Alloys“, ASM (1999) [7] Senf, J.: Untersuchung und Beschreibung von Magnesiumdruckgusslegierungen unter tribologischer, korrosiver und mechanisch-korrosiver Beanspruchung, Berichte aus der Werkstofftechnik, Shaker Verlag 2001, ISBN 3-8265-8428-7

632

Corrosion Stress Relaxation in Magnesium Alloys Ya. Unigovski, L. Riber, A. Eliezer, E. M. Gutman Ben-Gurion University of the Negev, Beer-Sheva, Israel

1

Introduction

The principal drawback of magnesium as a structural material, besides its relatively low corrosion resistance, is poor stress relaxation characteristics [1], which limit Mg alloys application in automotive and aerospace industry. It is well known [2, 3] that the relaxation behavior of a material is characterized by a spectrum of relaxation times or a relaxation spectrum H(τ). To the first approximation, H(τ)= - [dE(t)/dlnt]t=τ where τ is the relaxation time, E(t) = σ(t)/ε0 is the relaxation modulus; t is the current time, σ(t) is the current relaxation stress changing from the initial value of stress σ0 (t = 0) to its final value (the remaining stress) σrem, ε0 is the initial (usually elastic) strain corresponding to the stress σ0. The relaxation time corresponding to the maximum of the function H(τ), as well as the remaining stress, is accepted as characteristics of stress relaxation [2, 3]. Earlier we have found that the relaxation time strongly depends on casting parameters and porosity. It varied in AZ91D alloy at 150 °C from 10 to 20 hours [4]. The present paper deals with the corrosion stress relaxation behavior of pure Mg and diecast Mg alloys in 3.5% NaCl (pH ≈ 5) and 0.1N Na2B4O7 solution with a stable pH 9.3.

2

Experimental methods

Flat specimens (the total length 115 mm, gauge sizes 10 x 4 x 42 mm3) of polycrystalline magnesium (99.9653 % Mg) produced by Dead Sea Magnesium Works (Israel) were cut out from a commercial ingot in the transverse direction. Round specimens (5.9 mm in diameter, gauge length 75 mm) of Mg alloys AZ91D (Mg - 8.4% Al, 0.85% Zn, 0.17% Mn), AM50 (Mg - 5.1% Al, 0.15% Zn, 0.57% Mn) and AS21 (Mg - 2.3% Al, 0.23% Mn, 1.10% Si) were produced on a die-cast machine with the locking force of 2000 kN. Standard mechanical properties of pure Mg and its alloys at room temperature were as follows: ultimate tensile strength (UTS) equal to 60, 225, 229 and 221 MPa, tensile yield strength (TYS) equal to 24, 170, 136 and 134 MPa, elongation to fracture equal to 5.6, 2.8, 11.7 and 9.7% for pure Mg, AZ91D, AM50 and AS21 alloys, respectively. A decrease in tensile yield strength defined in air at 70 °C as compared to room temperature was equal to about 12±1 % for each alloy. Stress relaxation tension tests with the duration of 50 and 100 hours were performed on Zwick-1445 machine (Zwick GmbH & Co, Germany) at the temperatures of 25 ± 2 °C and 70 ± 1 °C and a constant strain during the entire test. The initial stress in stress relaxa-

633

Stress, %

tion tests varied from 0.67 to 1.15 of tensile yield strength of the alloys. This machine was completed by a split furnace Carbolite VST 12/75/300 and an electrochemical cell. The same cell was used earlier in corrosion creep tests [5]. Specimens were studied after rinsing with distilled water and wiping with acetone. The compositions of electrolytes for corrosion stress relaxation (CSR) tests were as follows: 0.1N Na2B4O7 buffer solution (pH 9.3) and 3.5 wt.% NaCl (pH ~ 5). Usually, CSR tests of pure Mg were carried out at a continuous solution feed at the rate of 3-5 cm3/min. Several sets of CSR tests of Mg alloys in the buffer solution were studied in the cell without the solution flow. During these long-term tests, some increase in the pH value was observed, but it was less than 3 %. All solutions were prepared from analytical grade chemicals and distilled water. All tests were carried out using at least 3 specimens for each set of experiments. Microstructure studies were carried out using an optical microscope “Nicon”, a scanning electron microscope JEOL JSM5600 and a transmission electron microscope JEM 2010. The size Figure 1: A crystallite in the cross-section (10 mm x 4 mm) of a of the magnesium crystallites is pure Mg sample. very large and equal to about 2 mm in diameter and more than 8 mm in the length (Fig. 1). The grain size of Mg-Al solid solution in the studied Mg alloys amounts to 1-10 microns and 10-15 microns in the surface layer and in the bulk of die castings, respectively. Some grains of AS21 alloy comprise also needles of silicide Mg2Si. Cleaning after tests implied the removal of 100 corrosion products from the surface of specimens by dissolution in 15% CrO3 solution at 80 80oC during 0.5-1 min (the ASTM standard G1-88). 60 40 20

3 A ir, % N a C l, %

Results and discussion

Relaxation tests were carried out for 50 and 100 hours. Since the remaining stress corre0 10 20 30 40 50 sponding to the end of tests was almost unT im e , h o u rs changed after the exposure exceeding 50 h, the majority of tests were performed for 50 h Figure 2: Stress relaxation in pure Mg in air and in 3.5%NaCl solution, 25 C, σ = 16 MPa (0.7 of (Table). TYS). Under the initial stress σ0 of 16 MPa or ~ 0.7 of tensile yield strength, pure magnesium shows a crucial decrease in stress in 3.5 % NaCl solution. For example, time-dependent stress after 50 hours decreases from 92.5 % in air to 21.2 % in a corrosive medium (Table, Fig. 2). At room temperature under the initial stresses σ0 varying in the range from ~0.8 to ~1.2 of tensile yield strength, all investigated alloys (AZ91D, AM50 and AS21) show a relatively small decrease in stress both in air and in different corrosive media (Table). For example, in AM50 alloy at σ0 = 115 MPa, the remaining stress σrem after 50 h amounts to 91.0%, 90.1% and 88.0% in air, 0.1N Na2B4O7 and 3.5% NaCl solutions, respectively 0

o

o

634 (Table). Growth of the initial stress from 115 MPa to 140 MPa leads to a decrease in σrem in AM50 from 91.0% to 81.0% in air, and from 88.0% to 80.1% in 3.5% NaCl solution (Table). Table. Effect of environment and the test temperature on stress relaxation in Mg alloys

Alloy

Medium

Temp. C

o

Initial stress, σ0/TYS σ0, MPa

25 150 25 195 70 115 70 150 25 150 0.1N Na2B4O7 3.5% NaCl 25 150*) 70 115 70 150*) 25 115 AM50 air 25 140 70 90 70 115 25 115 0.1N Na2B4O7 25 115 3.5% NaCl 25 140 70 90 70 115 air AS21 25 115 25 140 70 90 0.1N Na2B4O7 25 115 25 115 3.5% NaCl 25 140 70 90 *) A part of specimens was broken under this stress. AZ91D

air

0.88 1.15 0.79 1.03 0.88 0.88 0.79 1.03 0.85 1.03 0.77 0.98 0.85 0.85 1.03 0.77 0.98 0.86 1.04 0.78 0.86 0.86 1.04 0.78

Remaining stress, % 50 h

100 h

93.3 84.4 74.1 66.1 92.9 90.0 64.4 52.0 91.0 81.0 78.4 60.1 90.1 88.0 80.1 73.5 58.1 92.0 79.9 86.4 90.2 93.3 79.6 66.4

82.0 88.4 88.7 -

An increase in the test temperature from 25 °C to 70 °C leads to a significantly stronger stress relaxation of all alloys in comparison with room temperature, especially in corrosive solutions (Table). For instance, in air under the initial relative stresses close to the yield strength (σ0/TYS ~ 1) corresponding to a given alloy and temperature, the remaining stress decreases from 86.7% to 66.1% in AZ91D and from 81.0% to 60.1% in AM50, respectively. In 3.5% NaCl solution under the same relative stress, σrem for AM50 decreases from 80.1 % to 58.1 % with growth of the test temperature from 25 °C to 70 °C (Table). It is of interest that in 3.5% NaCl at the same absolute initial stress value of 115 MPa, AM50 and AZ91D alloys show almost the same relaxation resistance: σrem equal to 58,1% and 64.4%, respectively.

635 8.E+03 Air

AZ91D o 70 C 150 MPa

NaCl

H(t), MPa

6.E+03 4.E+03 2.E+03 0 .E+00 0

10

20 30 Time, hours

40

50

(a) 1 .E+04

8.E+03

H(t), MPa

Air

AM 50 o 70 C 90 MPa

NaCl

4.E+03

0.E+00

0

10

20 30 Time, hours

40

50

(b) 9 .E+03

Air NaCl

AS 21

H(t), MPa

o

70 C 90 MPa

6.E+03

3.E+03

0 .E+00 0

10

20 30 Time, hours

40

50

(c) Figure 3: Relaxation spectra of die-cast Mg alloys AZ91D (a), AM50 (b) and AS21(c) in air and in 3.5% NaCl solution at o 70 C under initial stress of 150 MPa (a) and 90 MPa (b, c).

At elevated temperature, the highest difference Δσrem between the remaining stresses corresponding to those in air and in 3.5% NaCl (a sensitivity to corrosive solution) was observed in AS21 alloy, and the lowest one – in AM50 alloy. For instance, in air and in 3.5% NaCl (70 °C, σ/ TYS ≅ 0.8), the remaining stress was equal to 74.1 and 64.4% (Δσrem = 9.7%), 78.4 and 73.5% (Δσrem = 4.9%), 86.4 and 66.4% (Δσrem = 20.0%) for AZ91D, AM50 and AS21 alloys, respectively (Table). However, in air at the same stress of 90 MPa, AS21 shows a higher relaxation resistance in comparison with AM50 alloy, namely, the remaining stress equals 86.4% and 78.4%, respectively. On the contrary, in 3.5% NaCl, AM50 manifests a higher relaxation resistance (remaining stress in AM50 and AS21 amounts to 73.5% and 66.4%, respectively). Thus, in a corrosive environment, the highest sensitivity to stress relaxation was observed both in the alloy with the highest Al content (AZ91D), as it was shown in creep tests [5], and also in AS21 alloy, which consists of two intermetallics (Mg17Al12 and Mg2Si). Calculated from the test data relaxation spectra show that AM50 alloy has very narrow peaks both in air and in corrosive environments. For instance, in Fig. 3 one can observe these peaks for all alloys in air and in NaCl at 70 °C. To the first approximation, the peak width or the maximum of the function H(τ) characterizes the material homogeneity. Usually,

636 the most homogeneous alloy gives a very narrow peak. Therefore, AM50 alloy having a very narrow peak shows the high stress relaxation homogeneity. On the contrary, AZ91D and AS21 manifest a significantly lower stress relaxation homogeneity (Fig. 3). The same behavior of AZ91D alloy was observed earlier in air at 150 °C, namely, under the initial stress of 0.8 of TYS, spectra of relaxation ti