PERGAMON MATERIALS SERIES VOLUME 2
Non-equilibrium Processing of Materials
PERGAMON
MATERIALS
SERIES
Series Editor: Robert W. Cahn F R S Department of Materials Science and Metallurgy, University of Cambridge, UK
Vol. 1 CALPHAD (Calculation of Phase Diagrams): A Comprehensive Guide
by N. Saunders and A. P. Miodownik Vol. 2 Non-equilibrium Processing of Materials edited by C. Suryanarayana A selection offorthcoming titles in this series:
Phase Transformations in Titanium- and Zirconium-based Alloys
by S. Banerjee and P. Mukhopadhyay Wettability at High Temperatures by N. Eustathopoulos, M. G. Nicholas and B. Drevet Ostwald Ripening by S. Marsh Nucleation by A. L. Greer and K. F. Kelton Underneath the Bragg Peaks: Structural Analysis of Complex Materials by T. Egami and S. J. L. Billinge The Coming of Materials Science by R. W. Cahn Thermally Activated Mechanisms in Crystal Plasticity by D. Caillard and J. L. Martin
PERGAMON MATERIALS SERIES
Non-equilibrium Processing of Materials edited by C. Suryanarayana The George S. Ansell Department of Metallurgical and Materials Engineering Colorado School of Mines Golden, Colorado, USA
1999
PERGAMON An Imprint of Elsevier Science Amsterdam- Lausanne- New Y o r k - Oxford- Shannon- Singapore- Tokyo
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Contents xv xvii xix
Series Preface Preface List of Contributors
CHAPTER 1 INTRODUCTION
1
R. W Cahn
CHAPTER 2 THERMODYNAMICS AND KINETICS OF METASTABLE PHASE FORMATION
5
K. N. Ishihara 2.1. Introduction 2.2. Thermodynamics of Metastable Phase Formation 2.2.1 Free Energy of Elements 2.2.2 Free Energy of Alloy Phases 2.2.3 Determination of Free Energy of Metastable Phases 2.2.4 Lattice Parameter of the Supersaturated Phase 2.3. Kinetics of Metastable Phase Formation 2.3.1 Nucleation of Metastable Phases 2.3.2 Nucleation of Alloy Phases 2.3.3 Crystal Growth Rate of Metastable Phases 2.4. Summary 2.5. List of Symbols References
5 5 5 7 9 11 12 13 15 18 19 19 20
CHAPTER 3 RAPID SOLIDIFICATION
23
H. Jones
3.1. Introduction 3.2. Methods of Rapid Solidification 3.2.1 Droplet Methods 3.2.2 Spinning Methods 3.2.3 Surface-Melting Methods 3.3. Constitution and Microstructure Formation by Rapid Solidification
V
23 24 24 26 27 28
3.3.1 Non-Equilibrium Constitution 3.3.2 Microstructure Formation 3.4. Properties, Performance and Applications of Rapidly Solidified Materials 3.4.1 Properties and Performance 3.4.2 Applications of Rapid Solidification References Selected Bibliography
CHAPTER 4 MECHANICAL ALLOYING
29 32 34 34 41 43 45
49
C. Suryanarayana 4.1. Introduction 4.2. Nomenclature 4.3. The Process of Mechanical Alloying 4.3.1 Raw Materials 4.3.2 Process Control Agents 4.3.3 Type of Mills 4.3.3.1 SPEX Shaker Mills 4.3.3.2 The Planetary Ball-Mill 4.3.3.3 Attritor Mills 4.3.3.4 Commercial Mills 4.3.3.5 New Designs 4.4. Mechanism of Alloying 4.5. Consolidation 4.6. Synthesis of Non-Equilibrium Phases 4.6.1 Solid Solubility Extensions 4.6.2 Synthesis of Intermetallics 4.6.3 Disordering of Ordered Intermetallics 4.6.4 Nanocrystalline Materials 4.6.5 Solid-state Amorphization 4.6.6 Displacement Reactions 4.7. Powder Contamination 4.8. Modeling 4.9. Industrial Applications 4.9.1 Nickel-Base Alloys 4.9.2 Iron-Base Alloys 4.9.3 Aluminum-Base Alloys 4.9.4 Magnesium-Base Alloys 4.10. Concluding Remarks References
vi
49
50 51 51 52 52
53 53 53 53 54 55 56 58 59 62 64 66 67 71 74 75 76 77 79 79 80 80 81
CHAPTER 5 LASER PROCESSING
89
K. F: Kobayashi 5.1. Principles of Lasers 5.2. Classifications of Laser Processing 5.3. Analysis of the Laser Melting and Quenching Process 5.3.1 Heat Transfer 5.3.1.1 Absorption Phenomena of Laser 5.3.1.2 Pulsed Laser 5.3.1.3 Continuous Wave (CW) Laser 5.3.2 Kinetic Conditions for Solidification 5.3.2.1 Kinetic Conditions for Amorphous Phase Formation 5.3.2.2 Solidification Modes in the Laser Processing of Materials 5.3.3 Microstructure Selection Maps (MSM) for Alloys 5.4. Laser-Quenching 5.4.1 Amorphous Phase Formation 5.4.2 Formation of Crystalline Phases 5.5. Laser Surface-Alloying and Cladding 5.5.1 Ferrous Alloys 5.5.2 Non-Ferrous Alloys 5.6. Laser-Annealing 5.7. Laser-Beam Joining 5.7.1 Welding of Structural Materials 5.7.1.1 Ferrous-Based Alloys 5.7.1.2 Non-Ferrous Alloys 5.7.2 Microjoining 5.8. Conclusions References
89 89 92 92 92 93 94 96 96 97 97 98 98 99 100 100 101 102 106 106 106 108 112 115 115
CHAPTER 6 THERMAL PLASMA PROCESSING
121
I? I!Ananthapadmanabhan and N. Venkatramani 6.1. Introduction 6.1.1 Advantages of Plasma Processing 6.2. Thermal Plasmas 6.2.1 Principles of Plasma Generation 6.2.1.1 DC Plasma Torches 6.2.1.2 AC Plasma Torches 6.2.1.3 F W Plasma Torches 6.2.2 Plasmagen Gases 6.2.3 Plasma-Particle Interaction 6.2.4 Plasma Processing Systems vii
121 121 123 123 124 126 126 127 128 129
6.3.
Processing of Materials 6.3.1 Plasma-Spraying 6.3.1.1 Structure of Sprayed Deposits 6.3.1.2 Plasma Spray-Deposited Materials 6.3.1.3 Low-Pressure Plasma-Spraying 6.3.1.4 Reactive Plasma-Spray Forming 6.3.1.5 Plasma Spheroidization 6.3.2 Plasma Reactors and Furnaces 6.3.2.1 Plasma Decomposition 6.3.2.2 Plasma Metallurgy 6.3.2.3 Processing of Ceramics 6.3.2.4 Treatment of Hazardous Wastes 6.3.3 Processing of Metastable Phases 6.3.3.1 Plasma Deposition of Diamonds 6.3.3.2 Thermal Plasma Synthesis of Ultrafine Alumina Powder 6.4. Summary and Conclusions Acknowledgments References
129 129 131 132 134 135 135 137 137 138 140 141 141 141 142 147 148 148
CHAPTER 7 153
SPRAY-FORMING Bing Li and E. J. Lavernia 7.1.
7.2. 7.3.
7.4.
7.5.
Introduction Principles Variations and Distinctions 7.3.1 Variations 7.3.1.1 Crucible: Induction Skull Melting/Spray-Forming 7.3.1.2 Atomizer: Circular vs. Linear Spray-Forming 7.3.1.3 Atomizer: Close-Coupled vs. Free-Fall Spray-Forming 7.3.1.4 Atomization Gas: Reactive Spray-Forming 7.3.1.5 Substrate: Near Net Shape Spray-Forming 7.3.1.6 Spray-Forming and Co-Injection 7.3.2 Nomenclature 7.3.2.1 Spray Atomization 7.3.2.2 Related Processing Techniques Applicability Non-Equilibrium Phenomena in Spray-Forming 7.5.1 Non-Equilibrium Nature 7.5.1.1 Rapid Solidification in Atomized Droplets 7.5.1.2 Transient Semi-Solid Layer in Deposition Stage 7.5.2 Non-Equilibrium Related Features in the Deposit 7.5.2.1 Metastable Phases 7.5.2.2 Extended Solid Solubility 7.5.2.3 Absence of Macrosegregation/Minimized Microsegregation viii
153 154 157 157 158 158 159 160 163 163 166 166 167 168 172 173 173 176 178 178 178 179
7.5.2.4 Refinement Effects of Non-Equilibrium Features on Mechanical/Physical Properties 7.5.3.1 Direct Effects 7.5.3.2 Indirect Effects 7.6. Concluding Remarks Acknowledgments References
180
7.5.3
181 183 185 189 189 189
CHAPTER 8 ION-MIXING
197
B. X. Liu
8.1. 8.2. 8.3.
Introduction Brief Description of Underlying Physics in Ion Mixing Thermodynamics of Alloy Phase Formation 8.3.1 Miedema's Theory and Alonso's Method 8.3.2 Interfacial Free Energy in the Multilayers 8.4. Experimentation of Ion-Mixing 8.4.1 Sample Design 8.4.2 Ion-Mixing Parameters 8.4.3 Characterization Methods 8.5. Amorphous Phase Formation 8.5.1 Glass-Forming Ability of Systems with a Negative Heat of Formation 8.5.2 Glass-Forming Ability of Systems with a Positive Heat of Formation 8.5.2.1 Amorphous Alloys Formed Within Restricted Compositions 8.5.2.2 Amorphous Alloys Formed in a Broad Composition Range 8.5.2.3 Nominal and Intrinsic Glass-Forming Ability 8.6. Formation of Metastable Crystalline Alloys 8.6.1 Structural Classification of the Metastable Crystalline Phases 8.6.1.1 Solid Solutions 8.6.1.2 h.c.p.-I and f.c.c.-I Phases 8.6.1.3 h.c.p.-II and f.c.c.-II Phases Based on b.c.c. Metals 8.6.2 Free Energy Calculation of the MX Phases 8.7. Interface-Generated Solid-State Vitrification in Systems with a Positive Heat of Formation 8.8. Concluding Remarks Acknowledgments References
197 199 200 200 201 203 203 205 205 205 206 207 207 207 211 213 213 213 213 215 216 218 219 220 220
CHAPTER 9 PHYSICAL VAPOR DEPOSITION
225
J. S. Colligon
ix
9.1. 9.2. 9.3.
Introduction Development of PVD Deposition Methods 9.3.1 Evaporation 9.3.2 Molecular Beam Epitaxy (MBE) 9.3.3 Sputtering 9.3.4 Ion-Assisted Deposition 9.3.5 Magnetron Sputtering 9.4. Influence of Energy on Coatings 9.4.1 Film Morphology and Density 9.4.2 Nucleation and Adhesion 9.4.3 Metastable Phases 9.4.4 Microstructure 9.5. Applications of PVD Coatings 9.5.1 Optical Coatings 9.5.2 Corrosion Protective Coatings 9.5.3 Hard Coatings 9.6. Future Trends Acknowledgment References
225 225 228 228 229 229 231 235 238 238 239 240 240 243 243 245 246 248 249 249
C H A P T E R 10 CHEMICAL VAPOR DEPOSITION
257
E Teyssandier and A. Dollet
10.1. Introduction 10.1.1 Presentation 10.2. Gas-Phase Transport and Reactivity 10.2.1 Non-Reactive Fluid Flow 10.2.1.1 Geometrical Effects 10.2.1.2 Influence of Operating Conditions 10.2.2 Reactive Flows 10.2.2.1 Thermodynamic Approach 10.2.2.2 Gas-Phase Mechanisms and Kinetics 10.2.2.3 Simulations of the Gas-Phase Composition 10.2.2.4 Nucleation and Growth of Solid Particles from the Gas Phase 10.3. Solid Phase Formation 10.3.1 Solid-Gas Thermodynamic Equilibrium Approach 10.3.2 The Structure of CVD Layers 10.3.3 The Driving Force for Crystal Growth from the Vapor Phase 10.3.4 Surface Mechanisms and Kinetics 10.3.4.1 Rate of Elementary Surface Processes 10.3.4.2 Estimation of the Kinetics of Surface Processes in CVD 10.3.5 Heterogeneous Nucleation 10.3.5.1 Nucleation Theories X
257 258 259 260 260 261 261 263 264 264 265 267 267 268 270 271 271 272 273 274
10.3.5.2 10.3.5.3 10.3.5.4 10.3.5.5 10.3.5.6 10.3.5.7
Two- and Three-Dimensional Nucleation Epitaxial Nucleation and Growth Multi-Component Nucleation and Chemical Reactions Nucleation Enhancement Selective Vapor Deposition Theoretical and Experimental Studies of Nucleation in CVD 10.3.6 Crystal Growth 10.3.6.1 Mechanisms of Crystal Growth 10.3.6.2 Classification of Crystal Faces 10.3.6.3 Crystalline Morphology 10.3.6.4 Morphology and Normal Growth Rate of Crystal Faces 10.3.6.5 Relation between Preferred Orientation and Morphology of Polycrystalline Films 10.3.6.6 Theoretical Studies of Crystal Growth from the Vapor Phase 10.4. Conclusions References
275 275 276 276 277 277 278 278 278 279 279 281 281 281 282
C H A P T E R 11 COMBUSTION
SYNTHESIS
289
S. B. Bhaduri and S. Bhaduri
11.1. Introduction 11.2. Thermodynamic Considerations 11.3. Kinetic Considerations 11.4. Field-Activated Combustion Synthesis 11.5. The "Azide" Process 11.6. SHS Reactions in Synthesizing Ti3SiC2 11.7. Controlled Reactions in the Ti-B Binary System 11.8. Auto-Ignition Synthesis of Nanocrystalline Oxides 11.9. Non-Equilibrium Effects 11.10. Concluding Remarks Acknowledgments References
289 291 294 295 297 299 301 304 307 307 308 308
C H A P T E R 12 NANOSTRUCTURED
MATERIALS
313
c. Suryanarayana and C. C. Koch
12.1. 12.2. 12.3.
Introduction Classification Preparation 12.3.1 Inert Gas Condensation
313 313 314 314 xi
12.3.2 Rapid Solidification 12.3.3 Electrodeposition 12.3.4 Chemical Reactions 12.3.5 Mechanical Attrition 12.3.6 Devitrification 12.4. Structure 12.4.1 Microstructure 12.4.2 Atomic Structure of the Crystal Lattice 12.4.3 Atomic Structure of the Grain Boundaries 12.4.4 Triple Junctions and Higher-Order Grain Junctions 12.5. Stability 12.5.1 Grain Growth in Nanocrystalline Materials 12.5.2 Grain Growth at Ambient Temperature 12.5.3 Examples of Grain Growth Inhibition 12.5.4 Isothermal Grain Growth Kinetics 12.6. Particulate Consolidation 12.7. Properties 12.7.1 Diffusion and Sinterability 12.7.2 Mechanical Properties 12.7.2.1 Elastic Properties 12.7.2.2 Hardness and Strength 12.7.2.3 Ductility and Toughness 12.7.2.4 Superplastic Behavior 12.7.2.5 Deformation Mechanisms in Nanoscale Materials 12.7.3 Magnetic Properties 12.7.3.1 Fundamental Properties 12.7.3.2 Soft Magnetic Materials 12.7.3.3 Hard Magnetic Materials 12.7.3.4 Other Ferromagnetic Nanocrystalline Materials 12.7.4 Chemical Properties 12.7.4.1 Corrosion Behavior 12.7.4.2 Catalytic Properties 12.8. Applications--Present and Potential 12.8.1 Structural Applications 12.8.1.1 Cutting Tools 12.8.1.2 Nanocomposites 12.8.1.3 Superplastic Materials 12.8.1.4 Coatings 12.8.2 Magnetic Applications 12.8.3 Catalysts and Hydrogen Storage Materials 12.8.4 Functional NanostructuresmElectronic Applications 12.9. Concluding Remarks References
xii
316 317 318 318 319 319 320 321 321 322 323 323 324 324 325 326 327 327 328 328 329 329 331 331 333 333 333 333 335 335 335 335 337 337 337 338 339 339 340 340 341 341 342
CHAPTER 13 POWDER CONSOLIDATION
347
J. R. Groza
13.1. 13.2. 13.3.
Introduction Metastability in Powder Consolidation Consolidation of Metastable Powders 13.3.1 Sintering Mechanisms 13.3.2 Scaling Laws 13.3.3 Powder Contamination 13.4. Consolidation Methods 13.4.1 Conventional Sintering 13.4.1.1 Cold Compaction 13.4.1.2 Warm Compaction 13.4.1.3 Conventional Sintering of Nanopowders 13.4.1.4 Nano-Composite Densification 13.4.1.5 Grain-Size Control 13.4.2 Pressure Consolidation of Metastable Powders 13.4.3 Pressure-Assisted Consolidation Methods 13.4.4 Non-Conventional Sintering Methods 13.4.4.1 Microwave Sintering 13.4.4.2 Field-Assisted Sintering 13.4.4.3 Shockwave Consolidation 13.5. Concluding Remarks Acknowledgments References
347 348 349 352 353 355 356 356 356 358 359 361 361 363 364 366 366 366 368 368 369 369
CHAPTER 14 BULK AMORPHOUS ALLOYS
375
A. Inoue
14.1. 14.2. 14.3. 14.4. 14.5. 14.6.
14.7. 14.8. 14.9.
History of Bulk Amorphous Alloys Dominant Factors for High Glass-Forming Ability Continuous Cooling Transformation of Alloys with High Glass-Forming Ability Preparation Methods and Elemental Effect Structural Relaxation and Glass Transition Physical Properties 14.6.1 Density 14.6.2 Electrical Resistivity 14.6.3 Thermal Expansion Coefficient Mechanical Properties Viscoelasticity Soft Magnetic Properties xiii
375 376 380 383 387 392 392 393 394 395 399 403
14.9.1 Formation and Soft Magnetic Properties of Bulk Amorphous Alloys 14.10. Viscous Flow and Micro-Formability of Supercooled Liquid 14.10.1 Feature of Phase Transition of Bulk Amorphous Alloys 14.10.2 Deformation Behavior of Supercooled Liquid 14.10.3 Micro-Forming of Supercooled Liquid 14.11. Applications and Future Prospects References
403 406 409 409 412 413 413
Author Index
417
Subject Index
419
xiv
Series Preface
My editorial objective in this new series is to present to the scientific public a collection of texts that satisfies one of two criteria: the systematic presentation of a specialised but important topic within materials science or engineering that has not previously (or recently) been the subject of full-length treatment and is in rapid development; or the systematic account of a broad theme in materials science or engineering. The books are not, in general, designed as undergraduate texts, but rather are intended for use at graduate level and by established research workers. However, teaching methods are in such rapid evolution that some of the books may well find use at an earlier stage in university education. I have long editorial experience both in covering the whole of a huge field- physical metallurgy or materials science and t e c h n o l o g y - and in arranging for specialised subsidiary topics to be presented in monographs. My intention is to apply the lessons learnt in more than 35 years of editing to the objectives stated above. Authors (and, in some instances, editors) have been invited for their up-to-date expertise and also for their ability to see their subjects in a wider perspective. I am grateful to Elsevier Science Ltd., who own the Pergamon Press imprint, and equally to my authors, for their confidence.
ROBERT W. CAHN, FRS
(Cambridge University, UK)
XV
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Preface
Metals and materials have been used for a long time. But the rapid technological developments during the later half of the twentieth century demanded materials that are stronger, capable of use at much higher temperature, more corrosion-resistant, and much less expensive than those currently in use. These demands become even more significant when we are on the threshold of the new century and the millenium. Significant improvements in properties can only be achieved by processing the materials under far-from-equilibrium (or non-equilibrium) conditions. Accordingly, several new processing technologies have been developed during the past few decades, including rapid solidification, spray-forming, mechanical alloying, ion-mixing, vapor-deposition, laser processing, plasma-processing, etc. New types of materials such as amorphous alloys and nanostructured materials have also been developed. Professor Robert Cahn has put these developments in a proper perspective in his introductory chapter, which also sets the tone for the chapters that follow. Remarkable advances have been made in recent years in the science and technology of these processes used to synthesize, characterize, and apply materials processed under non-equilibrium conditions. There have been many success stories. Some of these techniques have evolved from laboratory curiosities to commercial-scale manufacturing in a few short years. In other cases, industrial necessity prompted development of the technology, and the science followed later. The very large number of technical papers published in these different areas have been summarized in reviews, monographs, or conference proceedings, but have almost always been devoted to a single processing technique. I felt that a book of this sort was needed to provide a critical and comparative study of the different processes in a single volume. The different chapters in this book have been written by people who are worldrecognized experts in their respective fields. Each chapter describes the principles, processing techniques, special features of the materials produced and their applications. An extensive list of references is provided at the end of each chapter which will facilitate location of additional information on specific aspects of any technique. Even though great care has been taken to provide coherency and uniformity amongst the different chapters, the individual style of the authors has not been completely sacrificed. After all, variety is the spice of life! The book is primarily intended for use by graduate-level students of materials science and engineering. However, scientists beginning their research careers in any of the processes covered in this book can use it to get a fairly good knowledge and overview of the state of the technique and its potential. Scientists who wish to enter this area of research can also use this book to get a comparative idea of the potential and limitations of the different non-equilibrium processing techniques.
xvii
I wish to thank all the contributing authors for their considerable efforts in summarizing the voluminous literature of their specialties, and their dedication, diligence, and patience in working with me (and tolerating the many editorial iterations). The encouragement provided by Robert Cahn (Cambridge, UK) and John Moore (Golden, CO, USA) is greatly appreciated. I would like to express my appreciation to Phil Mestecky and his staff at Elsevier for bringing this book out in a timely and attractive fashion. Finally, I wish to thank my wife Meena for all her support during the preparation of this volume.
Golden, CO, USA August 31, 1998
C. SURYANARAYANA
xviii
List of Contributors
E V. ANANTHAPADMANABHAN Laser and Plasma Technology Division, Bhabha Atomic Research Centre, Mumbai400 085, India S. BHADURI Department of Metallurgical and Mining Engineering, University of Idaho, Moscow, ID 83844-3024, U.S.A. S. B. BHADURI Department of Metallurgical and Mining Engineering, University of Idaho, Moscow, ID 83844-3024, U.S.A. R. W. CAHN Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ, U.K. J. S. COLLIGON Manchester Metropolitan University, Chester Street, Manchester M1 5GD, U.K. A. DOLLET Institut de Science et de Grnie des Matrriaux et Procrdrs, CNRS, UPR8521, Universitr, Avenue de Villeneuve, F-66860 Perpignan Cedex, France J. R. GROZA Department of Chemical Engineering and Materials Science, University of California at Davis, One Shields Ave., Davis, CA 95616, U.S.A. A. INOUE Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan K. N. ISHIHARA Department of Socio-Environmental Energy Science, Kyoto University, Kyoto 6068501, Japan H. JONES Department of Engineering Materials, University of Sheffield, Mappin Street, Sheffield S1 3JD, U.K.
xix
K. E KOBAYASHI Department of Manufacturing Science, Graduate School of Engineering, Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan C. C. KOCH Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC 27695-7907, U.S.A. E. J. LAVERNIA Department of Chemical & Biochemical Engineering and Materials Science, University of California, Irvine, CA 92697-2575, U.S.A. B. LI
Department of Chemical & Biochemical Engineering and Materials Science, University of California, Irvine, CA 92697-2575, U.S.A. B. X. LIU
Department of Materials Science and Engineering, Tsinghua University, Beijing 100084, China C. SURYANARAYANA The George S. Ansell Department of Metallurgical and Materials Engineering, Colorado School of Mines, Golden, CO 80401-1887, USA E TEYSSANDIER Institut de Science et de G6nie des Mat6riaux et Proc6d6s, CNRS, UPR8521, Universit6, Avenue de Villeneuve, F-66860 Perpignan Cedex, France N. VENKATRAMANI Laser and Plasma Technology Division, Bhabha Atomic Research Centre, Mumbai400 085, India
XX
Chapter 1 Introduction ROBERT W. CAHN
Man has used some of his materials in non-equilibrium form for a long time: The Japanese samurai sword of the thirteenth century, with its martensitic edge; age-hardened aircraft alloys, introduced early in this century; vapor-quenched superconducting alloys in the 1930s; melt-spun transformer laminations consisting of metallic glass in the 1970s. It was the chance discovery of age-hardening by Wilm at the start of this century, supplemented by Merica's recognition that the hardening stemmed in some way from a supersaturated solid solution, that really drew metallurgists' attention to the practical importance of non-equilibrium states. From the mid-1930s onwards, crystallographers like Guinier and Preston, physical chemists like Turnbull and physical metallurgists like Nutting, devoted their skills to unravelling the nature of age-hardening, at the same time as other metallurgists such as Kurdyumov in Russia and Cohen in the USA began the long task of understanding the details of the martensitic transformation in steels. Metastability had become a way of life for some metallurgists. The attitude of metallurgists to metastability underwent a sudden sea-change when Pol Duwez in California, in 1960, invented what at first came to be called splat-quenching (a term which, incidentally, he hated) and later, rapid solidification processing. This novel technique, centred on rapid quenching of hot molten alloys, introduced to the materials community things like hugely supersaturated solid solutions and, particularly, metallic glasses. Like all important innovations, this one did not spring fully formed from its progenitor's fertile imagination. Some years earlier, Duwez had studied the cooling rates accessible with a thin solid foil and a fast-moving stream of helium, and found them limited to around 10,000 K/s. A number of engineers had used rapid freezing, for instance of wires, purely as a means of manufacturing convenient shapes cheaply. Duwez realized that, to make progress, he had to quench, not a solid but a melt. What was new about Duwez's researches was that, for him, rapid freezing was a means to enhance accessible cooling rates by a large factor, and thereby to generate new states of matter: it was a scientific study in the first instance, and it led to technological advances as a by-product. What went on at California Institute of Technology at the beginning of the 1960s was a textbook illustration of the principle that technological innovation often (by no means always) springs from new scientific insights, and from experiments designed to create such insights.
2
R.W. Cahn
Once it had become clear that strength, in particular, could be greatly enhanced by supersaturating solid solutions and then heat-treating such non-equilibrium structures, metallurgists began to examine other ways of achieving such structures. Mechanical alloying began in the nickel industry as a way of making useful nickel-base alloys, and was then adopted by academically inclined researchers as a way of producing alloys in non-equilibrium states: in particular, alloys such as Ni3A1 which in equilibrium are fully ordered could be produced in disordered form, and some of these artificially disordered alloys then turned into glasses; this was a kind of two-stage metastability. Physical vapor deposition, which as mentioned above had been studied since the 1930s, was looked at far more systematically, and new devices such as lasers (that notorious family of solutions looking for a problem) were pressed into service both for this and for other processing strategies. The microelectronic industry put great effort into chemical vapor deposition and into ion implantation and ion mixing, and some of this work gave new ideas to classical metallurgists. Some years ago, an influential committee of the National Academies of Science and Engineering in the USA pointed to the overwhelming importance of improvements in processing to the national economy. Non-equilibrium processing of the types mentioned can be said to have responded to this need. Research was not limited to the academic (as in much recent work on mechanical alloying and milling) or to the tiny, as in microelectronics. New major processing techniques were improved: plasma-processing and spray-forming are large-scale techniques which happen to incorporate non-equilibrium states of matter. Another approach, mostly pioneered in Russia, was combustion synthesis (which, like splat-quenching, passed through several names); long ago, this approach was used to weld together streetcar rails in situ, but it has come a long way since then, and recently has turned full circle as a way of improving casting technique. Finally, and in some ways the most exciting, is the development of research on nanostructures, introduced to the world of modem science by Gleiter in Germany. Nanostructural features affect both functional materials (quantum dots, soft magnets) and structural materials. A century earlier, something similar had formed the subject-matter of the science of colloid science, but this never really took off as a wholly respectable science: maybe, in part, this was because early colloid scientists had poured cold water on the hypothesis that giant (macro-) molecules can form from organic monomers, and held back the development of polymer science by a couple of decades at least. It is always hard to forgive willful obscurantists! Now, many years later, colloid scientists (they still exist) have made fascinating discoveries about "colloidal crystals", consisting of equisized colloidal particles rather than of atoms. There have been books devoted to rapid solidification, to nanostructures, to metallic glasses (before they found application in that most cogent of industries, that of golf), to laser-processing and to other techniques toombut this, I do believe, is the first book to cover these variegated techniques in a critical and comparative way. I have long known Prof. Suryanarayana as a masterly editor, and here he has excelled himself.
Chapter 2 Thermodynamics and Kinetics of Metastable Phase Formation 2.1. 2.2.
Introduction Thermodynamics of Metastable Phase Formation 2.2.1 Free Energy of Elements 2.2.2 Free Energy of Alloy Phases 2.2.3 Determination of Free Energy of Metastable Phases 2.2.4 Lattice Parameter of the Supersaturated Phase 2.3. Kinetics of Metastable Phase Formation 2.3.1 Nucleation of Metastable Phases 2.3.2 Nucleation of Alloy Phases 2.3.3 Crystal Growth Rate of Metastable Phases 2.4. Summary 2.5. List of Symbols References
5 5 5 7 9 11 12 13 15 18 19 19 20
This Page Intentionally Left Blank
Chapter 2 Thermodynamics and Kinetics of Metastable Phase Formation KEIICHI N. ISHIHARA
2.1. I N T R O D U C T I O N
Metastable phases are commonly seen in nature. Typical examples are the occurrence of martensite in steel and diamond form of carbon. Recently, non-equilibrium processes such as rapid solidification, mechanical alloying and vapor deposition have been developed enabling formation of a variety of metastable phases. The fundamentals of some of these processing techniques have been reviewed (Turnbull 1981, Shingu 1992, Boettinger and Perepezko 1993, Herlach 1994, Baricco 1995). In this chapter, the thermodynamics and kinetics of metastable phase formation will be treated. The thermodynamics of competitive phases will be first discussed followed by the nucleation problem of metastable phase formation. Thermodynamically, reaction or phase transformation can occur only when the free energy of the system decreases. However, there is the possibility that it may not occur. For example, the free energy of diamond is higher than that of graphite at ambient pressure and room temperature. But diamond is quite "stable" at room temperature and pressure without any noticeable change occurring. Put differently, thermodynamics tells us that a reaction cannot occur when the free energy of the system increases. In the following paragraphs, we will show how the formation of a metastable phase can be expected to occur. Then we will describe how a phase can be formed under a wide range of thermodynamical possibilities. This is kinetics in contrast to thermodynamics, which only tells us whether, or not, a phase can form. In this part, the classical nucleation theory will play an important role. Using this theory, Ostwald's step rule will show that it is the metastable and not the stable phase that is transformed.
2.2. T H E R M O D Y N A M I C S O F M E T A S T A B L E P H A S E F O R M A T I O N
2.2.1 Free energy of elements Figure 2.1 shows schematically the free energies of the or,/~, y and 3 solid phases and the liquid phase of a hypothetical element. In this case, the/3, V and 6 solid phases are metastable with respect to the ot phase since the free energies of these phases are higher
6
K.N. Ishihara
0
6 solid
2 solid solid Liquid
0
Tg
T~m
T~
Tm
Temperature Figure 2.1. Schematic free energy hierarchies.
than that of the u phase at all temperatures. Tm ~ represents the melting point of the ot phase, where the free energy of the u phase is the same as that of the liquid phase. Similarly, Tm ~ and Tmy represent the melting temperatures of the fl and y phases, respectively. The Gibbs free energies of the different competing phases can be compared, at fixed temperature and pressure, to determine the relative stabilities of the phases under equilibrium. At a temperature higher than Tm ~ the free energy of the liquid phase is the lowest, and those of the c~,/3, y and 6 phases are higher in that order. This means that while liquid is the most stable phase under these conditions, the or, /3, y and 6 phases are less stable. We call this the hierarchy of stability. The hierarchy of stability of the different phases changes through a crossover of the free energy curves during a change in temperature and/or pressure. We will, however, ignore the pressure effect in this chapter. In this case, for the range of undercooling from Tm ~ to Tm ~, ot is the stable phase, the free energy of which is less than that of the liquid phase. Only the u phase can be formed from the undercooled liquid in this temperature region. Below Tm ~ it is possible to solidify the fl phase directly from the liquid. Therefore, there are two different ways in which the liquid can solidify. Thermodynamics does not tell us which phase will form preferentially from the liquid state. This will be discussed in the section dealing with the kinetics. In Fig. 2.1, the free energy of the 6 phase is higher than that of the liquid phase at all temperatures. Therefore, the 6 phase cannot form from the undercooled liquid. However, it may form from the vapor phase, depending on the relative position of the free energy of the vapor and 3 phases. The position of the amorphous phase is also indicated in this figure. The amorphous phase is the frozen state of an undercooled liquid. Since the free energy of the amorphous
Thermodynamics and Kinetics of Metastable Phase Formation phase depends on the frozen temperature, called the glass transition temperature Tg, the free energy cannot be precisely determined. The free energy of the amorphous phase is always higher than that of the liquid phase. Experimental determination of the free energy hierarchy has been performed for various elements. When the difference of specific heat between the liquid and solid phases is known, the difference in free energy, A G can be calculated using the equation:
A G - - AHf(Tm Tm - T) - fT Tm ACpdT + T fTTm ACP T dT
(2.1)
where AHu is the enthalpy of fusion, Tm is the melting temperature, and ACp is the difference of specific heat between the liquid and solid phases as a function of temperature. If no experimental data are available, the best approximation is to set A Cp = 0. The various expressions for this can be found in Kelton (1991). The free energies of different metallic elements are summarized by SGTE (Scientific Group Thermodata Europe) and can be found in CALPHAD (1991, 15, 317).
2.2.2 Free energy of alloy phases The minimum degree of undercooling which is thermodynamically necessary for the diffusionless transformation of a liquid alloy into a solid solution for a given alloy composition, is expressed by the To curve in the phase diagram. Figure 2.2 shows part of a hypothetical binary phase diagram and the free energy vs. composition curves at a temperature 7"1. Since the common tangent from A to B is lower than the free energy of either the single phase liquid or solid phases, the most stable state at the concentration from CA to C B is a mixture of the liquid and solid phases. This is indicated as the two-phase region in the phase diagram. However, in the case of partitionless solidification that is often observed during rapid solidification, the position of the free energy curve becomes important. Thus, at 7'1, in the concentration range below Co, the solid phase is more stable than the liquid phase and above Co the liquid phase is more stable than the solid phase. The trace of Co at which the solid and liquid phases have the same free energy can be expressed in the phase diagram; it is called To curve indicated by the dashed line in Fig. 2.2. If the To curve plunges to very low temperatures, single-phase crystals cannot form from the melt. In this case, phase separation from the liquid phase depresses the solidification temperature to the glass transition temperature, Tg, where the liquid transforms to the amorphous phase. The system with the To curves, only slightly depressed below the stable liquidus, is a good candidate for partitionless solidification in the entire composition range (Herlach 1994). It is very important to know the position of the To curves in the phase diagram. Since the equilibrium phase diagram is a way of expressing thermodynamic data, there might be a simpler way of finding the approximate position of the To curve from the general features of the solidus and liquidus lines. The simplest guess, which often works, is to draw the To curve connecting the midpoints between the liquidus and solidus lines at a given temperature.
K. N. Ishihara
uid
~
0
0
i i i i
I
CA
CO
CB
i
(D .4,..a
c~ ~D T1
iquidus Solidus ~
Concentration Figure 2.2. Part of a hypothetical binary phase diagram and the free energy curves of the solid and liquid phases at temperature T 1.
If we know the free energy as a function of temperature and composition, the metastable phase diagrams including To curves and extrapolation of phase boundaries can be drawn. The free energy function of the binary system is now available for the purpose of calculating the phase diagram. This calculation of the phase diagram using the free energy function is called the CALPHAD (CALculating the PHAse Diagram) method (Saunders and Miodownik 1998). Thermocalc program module that contains the database of these functions is useful for this purpose. The recent edition of the ASM handbook on phase diagrams includes reference to the calculation of the phase diagrams (Massalski et al. 1990). The specific journals available for the assessment of free energy functions in connection with phase diagrams are "CALPHAD" and Journal of Phase Equilibria.
Thermodynamics and Kinetics of Metastable Phase Formation
9
2.2.3 Determination o f free energy o f metastable phases
Direct measurement of thermodynamic data of metastable phases is difficult and only a limited data is available. This has been done by measuring either the heat released during transformation of the metastable phase to the stable form or by measuring the electromotive force (emf). Measurement of the heat released on transforming to the stable phase is one of the simplest methods. Usually this measurement has been performed by DSC (Differential Scanning Calorimetry) (Michaelsen et al. 1997), especially to determine the heats of crystallization of the amorphous phases. The specific heat can also be measured by DSC. The specific heat and enthalpy of transformation for a metastable solid phase have been measured and the free energy is calculated using Eq. (2.1) for the pure metal gallium. An example of the experimentally determined free-energy hierarchy for gallium is shown in Fig. 2.3 (Bosio et al. 1973). ot is the stable phase and fl and y are metastable phases in Fig. 2.3. Measurements of emf for the stable and metastable phases have also been reported (Ishihara et al. 1993, G~xtner et al. 1996). In this method, only the partial quantity of free energy, activity, can be directly measured. Such a measurement for the Ag-Cu system is shown in Fig. 2.4. When any thermodynamic value is not available, the enthalpy of mixing based on Miedema's macroscopic atom model can be used (de Boer et al. 1988). This approach gives us not only the enthalpies of mixing for the liquid and solid phases for a binary alloy, but also the interfacial energy between solid phases. Since in the macroscopic atom model, atoms are conceived as 'blocks' of the element, assuming the size of the block as the grain size the enthalpy of the structure including the grain boundaries can be obtained. The enthalpy of nanostructures, frequently obtained by mechanical alloying and other
a 1500
~,
-
~ 10oo 500
0
100
200
300N~
T e m p e r a t u r e (K) Figure 2.3. Experimentallydeterminedfree energyhierarchyfor Gallium (Bosioet aL 1973).
10
K. N. I s h i h a r a
10000
.
.
.
.
.
.
.
.
.
I 9
1
@
1000 I:- ~
't
0 1O0
> r @
10
1
0.0
0.2
0.4
0.6
0.8
1.0
Cu Figure 2.4. Activity of Cu in Ag-Cu solid solution phase measured by the emf method. Solid circles and diamonds show the values measured using the CuSO4 solution and the solid electrolyte RbCulCI, respectively.
(a)
(b)
~xo
9
A
B
Concentration
A
B
Concentration
Figure 2.5. Changes of enthalpy by grain refinement for the case of systems with (a)
positive heat of mixing and (b) negative heat of mixing.
Thermodynamics and Kinetics of Metastable Phase Formation
11
methods, can be evaluated by this method. In this case, the enthalpy of the nanostructure should be close to that of the liquid or amorphous state so that the grain size reaches nanometer dimensions (Bakker 1998). The enthalpy for the simple mixing of powders can be indicated by a horizontal line in Fig. 2.5. When the grain size becomes small, the total enthalpy of grain boundary, AHgb changes to closely resemble the enthalpy of mixing of the liquid phase. In the case of a system which has a positive heat of mixing, A Hgb should raise its value, while in the case of the negative heat of mixing, A Hgb should increase at both low and high concentration ranges but may decrease in the middle concentration range, as shown in Fig. 2.5 (a) and (b). It means that during mechanical alloying, the enthalpy of mixing of the solid phase does not always increase and approach that of the liquid. 2.2.4 Lattice parameter o f the supersaturated phase
The lattice parameter of a solid solution phase varies linearly with composition; this is known as Vegard's law. If the crystal structures of the terminal solid solutions are different, the average atomic volume varies linearly with the concentration. A positive deviation from Vegard's law is often seen in nature (Najafabadi et al. 1993). Extension of solid solubility is commonly observed during non-equilibrium processing and the deviation from Vegard's law has been confirmed for many systems and for wide composition ranges. Several attempts have been made to explain this deviation on the basis of the theory of elasticity (Gschneidner and Vineyard 1962). However, this deviation is attributable to the excess energy of the solid and can be thermodynamically explained using the enthalpy of mixing. Here, a simple idea will be presented. One must first consider how much energy can be stored by the deviation from the ideal volume which is shown by the Vegard's law. This energy E can be expressed as (Ziman 1972): E = (B/v)(AV/V)
(2.2)
where B is the bulk modulus, 7/is the Griineisen constant and A V is the change in the unit cell volume. In the case of a system involving metals with a cubic structure, A V can be expressed in terms of the lattice parameter a, as A V~ V ~- 3 A a / a
(2.3)
where V and a are the ideal unit cell volume and lattice parameter, respectively, according to Vegard's law and Aa is the deviation of the lattice parameter from the ideal value. This energy should be equal to the excess enthalpy of mixing. In the case of a regular solution model, it can be shown that E
= CACB~']
(2.4)
where f2 is the interaction parameter. Then, A a / a = (v/3B)CACB~2
(2.5)
K. N. Ishihara
12 0.006
!
!
!
i
0.004
<~ 0.002
0.000
,
0.0
I
0.2
,
I
I
0.4
0.6
,
I
0.8
1.0
Cu Figure 2.6. Deviation of the lattice parameter form Vegard's law, Aa for the AgCu system. The open and filled circles show the reported results of ball-milled and rapidly-quenched samples, respectively (Najafabadi et al. 1993).
In this model, the deviation from Vegard's law for the regular solution model is parabolic, which is consistent with the reported experimental data. Moreover, the sign of the deviation is the same as the sign of the heat of mixing. This is the reason why a positive deviation from Vegard's law is often seen in nature. This is because, in the case of the system that has a negative heat of mixing, intermetallic compounds are formed rather than a supersaturated solid solution. In the Ag-Cu system, the calculated value of Aa calculated using the heat of mixing by the CALPHAD method (Murray 1984) matches moderately well with the experimental results shown in Fig. 2.6 for y = 2 and B calculated as the average of 1.0 x 10 ll N/m 2 for Ag and 1.5 x 10 ll N/m 2 for Cu.
2.3. K I N E T I C S OF M E T A S T A B L E P H A S E F O R M A T I O N
Many metastable phases are found in nature and have also been produced by the recently developed novel materials processing techniques. According to thermodynamics, as mentioned above, there is a possibility for both stable and metastable phases to form. So, which is more probable? This question can be answered by the kinetics of the process. There are two main reasons why metastable phases form. The first reason is that a metastable phase can nucleate before the stable phase can nucleate. The second reason is that the crystal growth rate of the metastable phase exceeds that of the stable phase. There have been attempts to draw kinetic phase diagrams (Chvoj 1983). However, no systematic study in this field has been performed as compared to the thermodynamic in-
Thermodynamics and Kinetics of Metastable Phase Formation
13
vestigations. In this section, the nucleation and the growth rates of the metastable and stable phases will be discussed.
2.3.1 Nucleation of metastable phases The free energy hierarchy, the relative magnitude of the driving forces for the possible phase transformations, does not tell us which phase is more likely to form under a given set of conditions. It is well known that according to the Ostwald step rule (Ostwald 1897) the first phase to appear will be the metastable phase which has a higher free energy than the stable phase. Here, the proof of the rule using the classical nucleation theory will be presented (Ishihara et al. 1985). The relative magnitude of the homogeneous nucleation frequencies for various phases from undercooled liquid metals and alloys is considered. The volume nucleation frequency, Iv, at a temperature T given by the classical nucleation theory can be expressed as follows:
Iv--Kvexp
-
AG*f(O)) kT
(2.6)
where Kv is a kinetic parameter, 0 is the wetting angle of the nucleus on the surface of a substrate material, and k is the Boltzmann's constant. A G* is the excess free energy of the critical nucleus and is expressed as: 16zr o'3l) 2 AG* -- -~- ( A G 2 )
(2.7)
where o" is the surface free energy between the liquid and solid, v is the molar volume, and A G is the free energy difference between the liquid and solid phases, o" may be empirically expressed as a function of the latent heat and molar volume as: o- -- 0.065AHfv -2/3
(2.8)
where A H f is the enthalpy of fusion. AG may be simply expressed by Eq. (2.1) when
ACp = O ,
as
AG = A H f ( T m - T)/Tm
(2.9)
where Tm is the melting temperature. From these equations, the following expression can be obtained: AG* = m167r(0.065)3AHf
3
T2
(2 10)
(Tin - T) 2"
The magnitude of Kv depends on the atomic transport rate in the liquid across the solidliquid interface. The value of Kv for the nucleation of the metastable and the stable phases may be considered equal since both the phases appear from the same liquid phase. So, in the case of homogeneous or fixed wetting angle nucleation, the nucleation frequency
K. N. Ishihara
14
expressed by Eq. (2.6) may be compared by calculating the differences between the value of A G* and A G .f (f refers to the metastable phase). The critical temperature, where the nucleation frequency of the metastable phase is equal to that of the stable phase, can be given as a solution to the equation AG* = AG .I
(2.11)
This equation has a significant solution under the condition that the ratio of enthalpies of fusion for stable and metastable phases, D, is less than unity:
D-
AH~f/AHf < 1.
(2.12)
When the above condition is satisfied, the solution of Eq. (2.11) is to --
A _ ~/-DA
(2.13)
Tr - x/~A where A is the ratio of the melting temperatures (T~m/Tm) and Tr is the reduced temperature (T/Tm). Since Ttm/Tm < 1, the critical temperature, to, should exist between Tm ~ and OK. A schematic illustration of AG* and A G .I vs. temperature is shown in Fig. 2.7. The condition (B < 1) should usually be satisfied because A Hf approaches A G at low temperatures and A G .I for the metastable phase is always less than that of the stable phase. Therefore, it may generally be expected that the nucleation frequency of the metastable phase is larger than that of the stable phase if the undercooling is large enough. Thus,
]! Meta~ table
o
I
0
to
A
Reduced Temperature (T/I'm) Figure 2.7. Example of the relationship of A G* for stable and metastable phases when A = 0.8 and B = 0.5. For the temperature region below t0, the nucleation of a metastable phase is more likely than that of the stable phase.
Thermodynamics and Kinetics o f Metastable Phase Formation
15
300
250 liquid --~ 3' o
200
[.~ 150 100
0
1
2
time (s) Figure 2.8. The trace of temperature during rapid cooling of Ga droplet. It is clearly shown that the sample is crystallized first into the metastable y phase (melting temperature 238 K). Then, it transforms into the metastable/3 phase (melting temperature 257 K) and finally it transforms to the stable ot phase. The free energies of these phases are shown in Fig. 2.3.
the Ostwald step rule tends to be valid for large undercooling or supersaturation. Experimental confirmation of this rule can be found in several alloy systems (Mackintosh et al. 1983, Levi et al. 1988, Cortella et al. 1993, Abe et al. 1994, Yoshioka et al. 1994, Yang et al. 1996). The sequence of phase changes during rapid cooling of Ga is shown in Fig. 2.8 (Mackintosh et al. 1983), confirming the Ostwald step rule.
2.3.2 Nucleation o f alloy phases
The question to be answered here is what would be the composition of the solid phase that nucleates first from the undercooled liquid of a given initial composition. Figure 2.9 shows schematically the possible nucleus composition at a given undercooling. For the undercooled liquid of composition Co, the formation of a solid phase with a composition between a and b reduces the free energy of the system. The magnitude of reduction in the free energy is the largest for the nucleus with the composition at c. But, the most probable nucleus composition can be explained by the relative value of the nucleation frequency. The theory of nucleation in a binary system has been presented earlier (Kamenetskaya 1967, Kelly and van der Sande 1982, Gressin et al. 1983, Ishihara, et al. 1985, Lin et al. 1988, Evans and Greer 1990, Thompson and Spaepen 1993, Ghosh 1994). According to Kamenetskaya (1967), when a solid nucleus of radius r with a concentration Cs forms
16
K. N. I s h i h a r a
e~t) Solid ~D
O Liquid
"""
i
Co
i
a Concentration
c
Figure 2.9. Schematic illustration of the free energy-concentration diagram explaining the magnitude of the driving force for nucleation of solid phase in a solid-solutionforming binary alloy system.
from a liquid of initial composition Co, the total free energy change of the system is given as"
A G --
1
3Nv
Gt(CI) + \ 3Nv
Gs(Cs) q
N
(2.14)
where N is the number of atoms in a system, Gl and Gs are the free energies of the liquid and solid phases and Ct is the liquid composition after nucleation. The mass conservation law gives the relation"
1
3--lqv- CI +
3Nv
(2.15)
The nucleus should have the composition that gives the smallest free energy increase on nucleation. Such a composition should satisfy the condition when the free energy surface shows a saddle point. Considering a binary system in which both the liquid and solid phases are ideal solutions, the surface energy can be expressed as: cr -- aA(1 -- Cs) -k- ffBCs
(2.16)
Thermodynamics and Kinetics of Metastable Phase Formation
1200
17
Co To
1000
800
e~
M 600 N
~' 400 To 3
200
O
0.0
0.2
I
i
i
0.4
0.6
0.8
P 1.0
Concentration Figure 2.10. Example of a calculated phase diagram with the ideal liquid and solid solutions. The curves SP and SQ show the most probable nucleus concentration for composition-independent surface energy and for composition-dependent surface energy, respectively.
where O"A and cr8 are the surface energies between the pure metals and the liquid. The results of calculation of the composition of the most probable nucleus as a function of temperature are shown in Fig. 2.10. This estimation agrees well with the experimental results (Fecht 1991). The area below MOSN, which is the trace of the compositions a and b in Fig. 2.9, refers to the region where the driving force for solidification from the initial liquid of composition Co is positive. The line SQ shows the trace of the saddle points. The temperature at which partitionless solidification is most probable, the intersection point of SQ and C = 0.2 in Fig. 2.10, is exactly To~3.
18
K. N. Ishihara
2.3.3 Crystal growth rate of metastable phases There have been only a few systematic studies of the growth rate. In general, the growth rate S can be expressed as a function of the undercooling by the equation (Jackson 1984)"
S-
So(AT) f
(2.17)
with So and fl (1 < fl < 2) as constants, and A T is the degree of undercooling. So and fi have no physical meaning. However, when fl of the metastable phase is larger than that of the stable phase, the growth rate of the metastable phase may become larger than that of the stable phase at a large undercooling, as shown in Fig. 2.11. Such a situation can occur when the stable phase has a faceted solid-liquid interface while the metastable phase has a molecularly rough interface (Tuckerman and Bechhoefer 1992). The experimentally investigated growth rates of the stable and metastable phase in the F e - C - X system are schematically shown (Magnin and Kurz 1988).
metastable phase
~D
o
stable phase
Tm(metastable) Tm(stable) Temperature Figure 2.11. Hypothetical relation of the growth rate for both stable and metastable phases. In this figure, fl of stable and metastable phases is 1.4 and 2.0, respectively. For SO,the value for the metastable phase is three times larger than that of the stable phase.
Thermodynamics and Kinetics of Metastable Phase Formation
19
2.4. S U M M A R Y
The thermodynamic nature of the metastable phases was discussed. The possibility of the formation of metastable phases was also discussed from the kinetic point of view-nucleation and growth of metastable phases. Furthermore, emphasis was laid on the relationship between the excess heat of mixing and the deviation of lattice parameter from Vegard's law.
2.5. L I S T O F S Y M B O L S
A a Aa B C
A Cp D E G AG A G*
AHf A Hgb
Iv Kv k N r S So T
Tm Tr V AV v fl y cr 0 f2
ratio of melting temperatures lattice constant difference of lattice constant bulk modulus concentration difference of specific heat ratio of enthalpy of fusion stored energy free energy difference of free energy excess free energy enthalpy of fusion enthalpy of grain boundary nucleation frequency kinetic parameter Boltzmann's constant number of atoms in a system radius of nucleus growth rate constant temperature melting temperature reduced temperature unit cell volume difference of unit cell volume molar volume constant Grtineisen constant surface energy wetting angle interaction parameter
20
K. N. I s h i h a r a
REFERENCES
Abe, T., Akiyama, A. and Onodera, H. (1994) ISIJ International, 34, 429. Bakker, H. (1998) Enthalpies in Alloys (TransTech Publications, Ztirich, Switzerland), p. 31. Baricco, M. (1995) Key Eng. Mater, 103, 1. Boettinger, W. J. and Perepezko, J. H. (1993) Rapidly Solidified Alloys: Processes, Structures, Properties, Applications, ed. Liebermann, H. H. (Marcel Dekker, New York), p. 17. Bosio, L., Cortes, R. and Defrain, A. (1973) J. Chem. Phys., 70, 357. Chvoj, Z. (1983) J. Non-Equilib. Thermodyn., 18, 201. Cortella, L., Vinet, B., Desr6, E J., Pasturel, A., Paxton, A. T. and van Schifgaarde, M. (1993) Phys. Rev. Lett., 70, 1469. de Boer, E R., Boom, R., Mattens, W. C. M., Miedema, A. R. and Niessen, A. K. (1988) Cohesion in Metals: Transition Metal Alloys (North-Holland, Amsterdam). Evans, E V. and Greer, A. L. (1990) in Principles of Solidification and Materials Processing, Vol. 2, eds. Trivedi, R., Sekhar, J. A. and Mazumder, J. (TransTech Publications, ZUrich, Switzerland), p. 741. Fecht, H. J. (1991) Acta Metall. Mater, 39, 1003. Gartner, E, Busch, R., Haasen, E and Bormann, R. (1996) in Metastable Phases and Microstructures, Vol. 400, eds. Bormann, R., Mazzone, G., Shull, R. D., Averback, R. S. and Ziolo, R. E (Mater. Res. Soc., Pittsburgh, PA), p. 119. Ghosh, G. (1994) Mater Sci. Eng., A189, 277. Gressin, E, Eustathopoulos, N. and Desr6, E (1983) Scripta Metall., 17, 711. Gschneidner Jr., K. A. and Vineyard, G. H. (1962) J. Appl. Phys., 33, 3444. Herlach, D. M. (1994) Mater Sci. Eng. Rept., R 12, 177. Ishihara, K. N., Maeda, M. and Shingu, P. H. (1985) Acta Metall., 33, 2113. Ishihara, K. N., Diaz de la Torre, S., Hosoe, T., Yamamoto, A. and Shingu, E H. (1993) in Processing Materials for Properties, eds. Heinen, H. and Oki, T. (TMS, Warrendale, PA), p. 667. Jackson, K. A. (1984) Mater Sci. Eng., 65, 7. Kamenetskaya, S. S. (1967) Sov. Phys. Crystallogr., 12, 71. Kelly, T. E and van der Sande, J.B. (1982) in Chemistry and Physics of Rapidly Solidified Materials, eds. Berkowitz, B. J. and Scattergood, R. O. (TMS, Warrendale, PA), p. 35. Kelton, K. E (1991 ) Solid State Phys., 45, 75. Levi, C. G., Jayaram, V., Valencia, J. J. and Mehrabian, R. (1988)J. Mater. Res., 3, 969. Lin, H. M., Kim, Y. W. and Kelly, T. E (1988) Acta Metall., 36, 2537. Mackintosh, A. J., Ishihara, K. N., and Shingu, P. H. (1983) Scripta Metall., 17, 1441. Magnin, E and Kurz, W. (1988) Metall. Trans., 19A, 1955, 1965. Massalski, T. B., Okamoto, H., Subramanian, E R. and Kacprzak, L. (1990) Binary Alloy Phase Diagrams (ASM Intemational, Materials Park, OH). Michaelsen, C., Gente, C. and Bormann, R. (1997) J. Mater. Res., 12, 1463. Murray, J. L. (1984) Metall. Trans., A15, 261. Najafabadi, R., Srolovitz, D. J., Ma, E. and Atzmon, M. (1993) J. Appl. Phys., 74, 3144. Ostwald, W. (1897) Z. Phys. Chem., 22, 22. Saunders, N. and Miodownik, A. P. (1998) CALPHAD: A Comprehensive Guide (Elsevier, Oxford). Shingu, P. H. (ed) (1992) Mechanical Alloying (TransTech Publications, Ztirich, Switzerland). Thompson, C. V. and Spaepen, E (1993) Acta Metall., 41, 2021. Tuckerman, L. S. and Bechhoefer J., (1992) Phys. Rev., A46, 3178. Tumbull, D. (1981) Metall. Trans., 12A, 695. Yang, G., Zhuang, H. and Biswas, P. (1996) NanoStructured Mater., 7, 675. Yoshioka, M., Hancock, B. C. and Zografi, G. (1994) J. Pharmac. Sci., 83, 1700. Ziman, J. M. (1972) Principles of the Theory of Solids, second edition (Cambridge University Press, Cambridge), p. 67.
Chapter 3 Rapid Solidification 3.1. 3.2.
Introduction Methods of Rapid Solidification 3.2.1 Droplet Methods 3.2.2 Spinning Methods 3.2.3 Surface-Melting Methods 3.3. Constitution and Microstructure Formation by Rapid Solidification 3.3.1 Non-Equilibrium Constitution 3.3.2 Microstructure Formation 3.4. Properties, Performance and Applications of Rapidly Solidified Materials 3.4.1 Properties and Performance 3.4.2 Applications of Rapid Solidification References Selected Bibliography
23 24 24 26 27 28 29 32 34 34 41 43 45
This Page Intentionally Left Blank
Chapter 3 Rapid Solidification H. JONES
3.1. I N T R O D U C T I O N
The term solidification simply means formation of solid so the parent material(s) can be gaseous (vapor) or liquid (e.g., melt). Thus, solidification can involve continuous or discontinuous deposition with or without a chemical reaction or passage of a solidification front through a volume of melt. The term rapid solidification is normally applied to solidification from the melt. The description "rapid" can be taken to imply a short-time interval between initiation and completion of solidification and a high velocity of propagation of the advancing solidification front. Such rapid solidification is most readily achieved by imposing a high cooling rate during solidification. This requires a thin layer, filament or droplet(s) of melt to be generated in good contact with an effective heat sink. Rapid solidification then occurs either directly, as a result of coupling between external heat extraction and the transport of latent and specific heat required to propagate the solidification front, or indirectly during the recalescence that follows nucleation of solidification at large undercooling. Rapid solidification processes such as melt-atomization, spray-deposition, splat-quenching, melt-spinning and planar flow-casting and their derivatives all generate high cooling rates during solidification and their products contain structural features formed at different levels of prior undercooling. More control over the nucleation stages of rapid solidification can be obtained by slow cooling of a volume of melt (which can be quite large) that has been freed of agents than would otherwise catalyze nucleation at low undercooling. Nucleation of rapid solidification can then be triggered by contact with a nucleating phase at predetermined levels of undercooling, though in this case the slow cool that follows completion of solidification may be insufficient to preserve to ambient temperature some of the non-equilibrium features generated by the rapid solidification step. Correspondingly, growth velocity during rapid solidification can be coupled to the velocity of a heat source (such as a laser or electron beam) traversing the surface of a block of material. In this case the heat source generates a stable melt-pool which melts new material at its leading edge as fast as previously melted material refreezes at its trailing edge. Melting and freezing front velocities increase from zero at the bottom of the melt-pool to a value related to the traverse velocity at the sample surface. The present short survey is in three sections. The first section reviews methods of rapid solidification from the melt and their characteristics. The second section reviews current knowledge of the effects of 23
H. Jones
24
rapid solidification on constitution and microstructure formation. The third section gives examples of resulting property changes and of current or potential applications.
3.2. M E T H O D S
OF R A P I D S O L I D I F I C A T I O N
These are most conveniently discussed under the headings droplet, spinning and surface melting technologies. Droplet technologies all involve generation of multiple or single droplets of melt which solidify either individually or in a state of partial or complete coalescence on a deposition surface. Spinning technologies involve stabilization of a flowing melt-stream so that it freezes as a continuous filament, ribbon or sheet in contact with a moving chill surface or heat extracting fluid. Surface melting technologies involve rapid melting at a surface followed by rapid freezing sustained by rapid heat extraction into the unmelted bulk. Cooling rates during solidification can reach 101~ K/s in cross sections as small as ~ 0 . 1 / z m with undercoolings that can attain hundreds of K, front velocities as high as 100 m/s and solidification times as short as nanoseconds.
3.2.1 Droplet methods These are all variants of the long-established technology of making lead shot by teeming molten lead through a preheated steel dish perforated by an array of equal sized holes (Johnson et al. 1976). The fundamental principle is that a pendant drop or free falling melt-stream tends to break up into droplets as a result of surface tension. This process can be intensified, for example, by impingement of high velocity jets of a second fluid, by centrifugal force at the tip of a rotating cup or disc or by an applied electric field. The standard method, known as atomization, uses high velocity gas or water jets to rapidly fragment a volume of melt into large numbers of small droplets (Fig. 3.1) (Schmitt 1979). The jets form, propel and cool the droplets that may freeze completely in flight to form powder or Metal
Metal
Connifed~ II1 Figure 3.1. Principle of spray droplet (atomization)by impingementof high velocity gas jets on to a free falling or emergent melt-stream (Schmitt 1979).
Rapid Solidification
25
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Figure 3.2. Particle size distribution for Sn-5wt%Pb alloy as a function of atomizing gas pressure (Anderson 1991).
may form individual splats or a thick deposit by impact on to a suitable substrate. Powder or shot particulate generated by such techniques displays a range of particle sizes and shapes even for a single set of operating conditions. Increasing atomizing gas pressure, for example, then increases the yield of finer particle sizes (Fig. 3.2) (Anderson 1991) while the more effective quench of water atomization leads to more irregular shapes. Although particles more or less identical in size and shape from the same powder sample can show quite different microstructures (because of different local conditions), smaller particles tend to cool more rapidly and/or undercool more prior to solidification so tend to solidify more rapidly (i.e., at higher front velocity). Splats formed from droplets of given size tend to solidify even more rapidly because of more effective heat extraction from the larger surface area they offer, especially when at least one of their surfaces is in good contact with an efficient heat sink, such as a water-cooled rotating copper drum. A spray deposit can maintain the same microstructure as the equivalent splats provided that their solidification time is sufficiently less than the time interval between deposition of successive splats at a given location on the substrate (Oguchi et al. 1990). Atomizers range in size from laboratory units with capacities of less than 1 kg per run to commercial scale facilities with capacities as large as 50,000 tonnes per year (Anon 1989). Powders destined for high performance applications tend to be atomized with inert gases or in vacuum to minimize formation of oxides or other potentially damaging inclusions. The variability of powder particle sizes, shapes and microstructures within bulk atomized powder samples, splats and spray deposits (which is acceptable for many purposes), and the need to consolidate the materials for typical engineering applications, has stimulated work on droplet methods that employ lower cooling rates and allow measurement or control of the temperature at which nucleation occurs. Such methods include emulsification techniques in which a distribution of droplets is produced within a carrier liquid which can be inert, which can act as a flux for impurities in the droplets that otherwise would activate nucleation at low undercoolings or which can trigger nucleation at a spe-
26
H. Jones
cific undercooling (Perepezko 1984). A variant uses rapid solidification to produce a fine secondary network of eutectic in a single- phase matrix. This network is then spheroidized by heat treatment above the solidus to produce a distribution of liquid droplets in the solid matrix (Prasad et al. 1984). Information regarding melting and nucleation temperatures of the droplets can then be obtained by controlled heating and cooling coupled with differential calorimetry, and the results related to microstructures formed in the solidified droplets (Kim and Cantor 1994). Undercoolings of hundreds of degrees can be generated in relatively large volumes of melt by prior removal of impurities e.g., by fluxing (Kui et al. 1984), by levitation (containerless) melting (Weber et al. 1991), or by induction remelting in contact with a water-cooled copper hearth. A recent triumph of this latter approach has been to form substantial pieces of zirconium-based metallic glasses that display ultra-high strengths and good fatigue characteristics along with extensive plasticity in the supercooled liquid range (Johnson 1996). Benefits of the related levitation melting approach have included much fundamental information on the formation of solidification microstructures at very high undercoolings and growth velocities that cannot be accessed by other rapid solidification techniques (Herlach 1994, Herlach et al. 1993). 3.2.2 Spinning methods
These derive from the simplest system in which a single melt-stream emerging from an orifice is stabilized by surface film formation or solidification before it can break up into droplets (so-called melt-extrusion or melt-spinning), to the most sophisticated in which
gas pressure
eit
heoting winding
./~
~ /orifice
/ribbon
rotot/ng drum Figure 3.3. Schematicof free-jet chill-block melt-spinning (Overshott 1979).
Rapid Solidification
27
one or more melt-streams are used to make wide or composite ribbons by impingement on single or twin chill roll surfaces. In standard free-jet chill-block melt-spinning (Fig. 3.3) (Overshott 1979) the melt-stream forms a single ribbon typically 10 to 50 # m thick and a few millimeters wide by impingement on to a single chill roll rotating at a surface speed of some tens of meters per second. In planar flow-casting the distance from the nozzle orifice to the chill surface is reduced to less than ~0.5 mm to eliminate formation of a melt-pool and associated instabilities. Thin strips up to 500 mm wide have been cast in tonnage quantities by this method, either for direct application or pulverization into platelet particulate for consolidation into other product forms. Melt-extraction (Maringer and Mobley 1974), in which the rotating chill roll forms the products by direct contact with the surface of a crucible of melt or the melted extremity of the solid electrode, and its derivative melt-overflow (Boulby and Wood 1986a, b), do not involve generating a melt-stream. Continuous rapidly solidified filament with round section can be produced by the rotating water bath process in which a melt-stream emerging from a cylindrical orifice is solidified directly on entry into a rotating annulus of water contained within the lip of a rotating water bath (Ohnaka 1985). This process can also be used to generate powder particulate when the conditions result in the break-up of the melt-stream on entry into the water bath (Raman et al. 1982). Free-jet chill-block melt-spinning and meltextraction can be used to generate a flake product directly if continuity is interrupted by a series of regularly-spaced notches on the chill-block. Melt-spun ribbon or planar flow cast sheet can be pulverised, for example, by means of a blade-cutter mill (Gelinas et al. 1988, Pelletier et al. 1990) into platelets (_~0.3 mm across). These provide a more suitable feedstock for consolidation by powder metallurgy than continuous ribbon, for example, which exhibits a much lower packing density.
3.2.3 Surface-melting methods These derive from spot or traverse welding techniques and differ from them only in that the depth melted is limited to ensure that the ensuing solidification will be sufficiently rapid. In its simplest form, a single pulse or continuous traversing heat source is used to rapidly melt the surface of a block material, the unmelted bulk acting as the heat sink during the subsequent rapid solidification (Fig. 3.4) (Lux and Hiller 1972). The resulting rapidly solidified material has the same composition as the underlying parent material, although the rapid solidification may produce a different microstructure and much improved properties. A second possibility is to preplace or inject alloy or dispersoid additions at its surface so they are incorporated into the melt-zone to form a surface region of composition different from the underlying bulk. The third possibility is to melt a different material preplaced on the surface so that mixing with the underlying material is limited to the minimum required for effective bonding. All three variants offer the practical possibility of generating a more durable surface on an underlying material that is in all other respects entirely adequate for the application in view. Both nanosecond and picosecond power laser sources have been used to generate some quite spectacular non-equilibrium effects in surface melt-zones as shallow as 0.1 /xm in which cooling rates during solidification have been estimated to reach 10 l~ K/s (von Allmen et al. 1984) or more and solidification times to be as short as 10 -9 s (Spaepen 1987). Both traversing laser and electron beams
28
H. J o n e s
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Figure
have been used to treat entire surfaces via repeated incremental lateral displacement of the beam by the width of the melt-zone at the start of each new traverse (Lux and Hiller 1972). Samples with crack-free, relatively smooth treated surfaces can now be produced by appropriate control of the process parameters. The technique can also be used to develop coupling between the traversing beam and the solidification front so that the effects on resulting microstructure of systematic variations in front velocity up to -~ 1 m/s can be determined experimentally (Boettinger et al. 1984, Zimmermann et al. 1989). (See also Chapter 5 on "Laser Processing" by K.E Kobayashi in this book.)
3.3. C O N S T I T U T I O N AND M I C R O S T R U C T U R E F O R M A T I O N BY RAPID SOLIDIFICATION
Rapid solidification produces changes in constitution because the large undercoolings and front velocities involved promote formation of non-equilibrium phases and extensions in composition range of surviving equilibrium phases. Microstructural differences include changes in mode of growth and size refinement as a result of the short diffusion distances imposed. These aspects will be dealt with separately.
29
Rapid Solidification
3.3.1 Non-equilibrium constitution Non-equilibrium phases formed by rapid solidification include glasses, quasicrystalline phases, new crystalline phases and equilibrium phases with extended composition ranges. Glass formation requires that the entire or residual melt is undercooled below the glass-forming temperature before the advance of crystal formation can complete the solidification process. Glass formation is highly sensitive to composition, with pure metallic melts in particular requiring very high cooling rates during solidification to suppress the rapid advance of crystal formation. Alloy systems showing deep eutectics, however, form glasses rather more readily in the composition ranges near to these eutectics. Critical cooling rates for glass formation can be predicted by modelling the nucleation and growth of the competing crystalline (or quasicrystalline) phases. The melt will become a glass if cooling is rapid enough not to intersect the knee of the time-temperature-transformation (TTT) curve of the competing crystal forming process (Fig. 3.5) (Davies et al. 1974). In the absence of such diagrams, easy or bulk glass formation has been variously associated with ratio of glass formation to melting temperatures, large intervals between glass formation and crystallization temperatures (the so-called "supercooled liquid" range), low thermodynamic driving force for crystallization, large negative heats of mixing, simultaneous precipitation of more than one phase, large atomic size differences between the components, as well as deep eutectic formation (Peker and Johnson 1993, Massalski et al. 1981, Inoue et al. 1993). (See also Chapter 14 on "Bulk Amorphous Alloys" by A. Inoue in 2000
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30
H. Jones 9
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this book.) Figure 3.6 (Alonso et al. 1988) shows a glass formability map for binary solute additions to nickel in which high enthalpy of formation arising from atomic size difference and large negative enthalpy of mixing are seen to correlate well with glass formation under melt-spinning conditions (cooling rate ~ 106 K/s). Formation of non-equilibrium crystalline phases by rapid solidification occurs when their nucleation and growth kinetics results in a higher rate of formation than the competing equilibrium phase. Formation of metastable (cementite) cast irons rather than graphitic cast irons at temperatures below the austenite-cementite eutectic temperature is a well-known example for relatively modest cooling rates. Possibilities for such metastable phase formation multiply manyfold at the high undercoolings and cooling rates that typify processes such as melt-spinning and splat quenching. The TTT diagram in Fig. 3.7 (Herlach et al. 1992) illustrates the competitive kinetics of formation of three contending intermetallic phases and or-A1 in undercooled Al-12 at %Mn alloy. Based on assumptions made, equilibrium A16Mn should form at cooling rates > 5 • 103 K/s, displaced by T-
Rapid Solidification
31
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Figure 3.8. Maximum reported extensions of solid solubility of various solutes in aluminum. Black bars indicate the increases beyond equilibrium solid solubilities achieved by rapid solidification (Midson and Jones 1982).
32
H. Jones
phase at 5 x 103 to 2 x 105 K/s and an icosahedral phase at 2 x 105 to 2 x 106 K/s, in accord with experimental results. Extension of solid solubility by rapid solidification, again, has been widely observed, and can amount to orders of magnitude in some cases where equilibrium solid solubilities are very low (Fig. 3.8) (Midson and Jones 1982). While limited atomic size and electrochemical differences between solvent and solute appear to be necessary conditions for such solubility extension, the actual outcome is expected to be determined by conditions at the advancing solidification front. Fig. 3.9 (Boettinger et al. 1987, Kurz et al. 1988) shows the continuous increase in solute concentration with increasing solidification front velocity for ot-Ag solid solution in the Ag-15 wt%Cu alloy. The experimental data match predictions of dendrite growth modeling and approach 15 wt%Cu (segregation-free solidification, SFS) at 0.4 m/s above which plane front solidification is expected to be stabilized. The phenomenon of solute trapping is also expected to take effect in this velocity range in which the time for incremental advance of the solidification front is less than the time required for diffusional jumping of solute atoms so that instead they are engulfed by the advancing solid. Fig. 3.10 (Boettinger et al. 1984) shows the corresponding effect of varying alloy composition as well as growth velocity. The increase in velocity required for SFS with increasing Cu alloying content up to 15 wt% can be associated with incidence of plane front solidification while the decrease in this velocity above 15 wt%Cu can be attributed to control by solute trapping (Jones 1988).
3.3.2 Microstructure formation Figure 3.10 (Boettinger et al. 1984) is also an example of a solidification microstructure selection diagram (SMSD) in which dominant solidification mode is mapped on a plot which has a material parameter, such as alloy composition, as one axis and a processing parameter, such as front velocity, as the other axis. Figure 3.11 (Boettinger 1982) shows a further example in which the competition is between glass formation, formation of primary Pd3Si, y and x phases and formation of the Pd3Si-K eutectic. Modelling of the competitive growth situation implicit in such diagrams presents a challenge for solidification theory. The fundamental premise is that the solidification mode with the highest growth temperature will dominate under given conditions. Figure 3.12 (Gilgien et al. 1995) shows a comparison between experimental fact and predictions of such competitive growth modelling for Al-rich A1-Fe alloys, which form the basis of several high performance rapidly solidified aluminium alloys. The level of agreement is encouraging. While growth temperature can be crucial in determining which constituent and which mode of solidification prevails, an equally important consideration for performance and properties is the size-scale of the resulting microstructural constituent. Measurements of grain size ~.0 vs. cooling rate T (Fig. 3.13) (Jones 1991), suggest a relationship of the form ~.0 - K l J ' - n with n -- 1 and K1 ~ 107 # m K/s for Ni-5wt%A1 and Ti6wt%A1-4wt%V and ~ 106 # m K/s for Fe-6.3wt%Si and type 316 stainless steel. A prediction for pure aluminium assuming homogeneous nucleation and linear growth kinetics gave n = 0.9 and K ~ 2 x 107 (K/s) n (Boswell and Chadwick 1977). The recent observations of nanoscale dispersions of extended and non-equilibrium phases often adjacent or related to glass-forming composition ranges has also been attributed to
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operation of homogeneous nucleation (Greer 1994). Spacings of cellular or-A1 in A1Fe and A1-Cu alloys grown at 0.03 to 1 m/s show good agreement with predictions of the tip stability model of Kurz et al. (1986) and a newer array model of Hunt and Lu (1996). An example is shown in Fig. 3.14 (Lu et al. 1994). The tip stability model predicts ,k = K 2 / V 1/2, where K2 is 4Jr(DF/ATo) 1/2, in which D is the solute diffusivity in the melt, 1-" is the ratio of interfacial energy to entropy of fusion and AT0 is the alloy freezing range. Approximate analytical models predict primary dendrite spacing )~1 = K3G1/2V 1/4 with K3 = K4(AToDI') 1/4, where G is temperature gradient, K4 is 2.83k 1/4 (Hunt 1979) or 4.3/k 1/4 (Kurz and Fisher 1981) with k as a solute partition coefficient, and experimental data lend some support to this relation (Moir and Jones 1991). Secondary dendrite arm spacing )~2 is predicted to follow )~2 -- K5 i?-n with K5 "~ 50/xm (K/s) n and n = 1/3, as exemplified for or-A1 in dilute A1-Cu and AI-Si alloys in Fig. 3.15, although the experimental data for other systems show 0.25 < n < 0.5 and 16 < K5 < 160 # m K n s -n (Jones 1984). Predictions for the interphase spacing of a regular eutectic give )~ = K 6 / V 1/2 (Jackson and Hunt 1996), where 7 < K6 < 50 # m 3/2 s 1/2 for Al-based eutectics (Juarez-Islas and Jones 1987). Figure 3.16 shows eutectic spacing A vs. V for the A1-Cu eutectic, for which K6 ~ 9.4 # m 2/3 s-1/2 right up to the velocity for extinction of stable eutectic separation at V ~ 0.5 rrds (Jones 1984, Zimmermann et al. 1989).
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!
25
~l 30
COMPOSITION (wt % Cu)
Figure 3.10. Conditions for microsegregation free solidification of a-Ag as a function of growth velocity and alloy composition for Ag-rich Ag-Cu alloys (Boettinger et al. 1984).
3.4. P R O P E R T I E S ~ P E R F O R M A N C E
AND A P P L I C A T I O N S O F R A P I D L Y S O L I D I F I E D
MATERIALS
These will be discussed under the separate headings of Properties and Performance, and Applications (Sections 3.4.1 and 3.4.2).
3.4.1 Properties and performance Metallic glasses can exhibit high flow stresses and hardness with fracture stress approaching E/50, the theoretical limit, where E is elastic modulus. This can be combined with ductile behavior in bending (Fig. 3.17) (Davis 1976), shear and compression. Young's and shear moduli are typically 20-40% lower than in the crystalline state because the atoms in the glass are not constrained by a lattice to make self-similar displacements. Bulk modulus is however only a few percent lower in the glassy state reflecting the dense random packing which gives a density only 1-2% less than for the crystalline state. Further increments in hardness and strength along with increased stiffness can be obtained by precipitation of nanocrystals within a metallic glass matrix so as to suppress the inhomogeneous shear process (Fig. 3.17) that characterizes flow and fracture of metallic glass.
35
Rapid Solidification ,
-
25 A
E
E
I0
>~-
5
m
1
I
J
~
I
o ~, ~ ~ \ 0
'
9 Gloss
\\
/
~o~._ "\ - ~ ~ \
N
~
I
t
i
i
'i
I
l
I 15
!
a
/
\',
%~
~
,/X
\~
-
_
i '
/
g
/
/
Gio,=+x z~
9
A
/k
~2.5
> ::E I'--
tw (..9
_
D Pd 3 s i + X + ~
"
/
eut~tic
///
0.5
0.25 -
"~' J
~ I
I
l
0
r 1 5
\~
\~
o //
\ \
l
/
K + euttctie
o
/
1 ~ ~ AT. % Cu
I l0
~
I
A
I
Figure 3.11. Constituent selection diagram for Pd-Cu-Si alloys for solidification front velocities of 0.25-50 mm/s (Boettinger 1982).
1000
AI plane front '
. . . . . . .
'
- ~,~,-,~ ,-,i ,-,2\",'~ . . . . . . . .
100
.=.
10
AI cell / dend. + AI-AIsFe eul.
E u
,~TITITIT11]Trr[~
L---,
>
I-l~Tr'-'-
r
/! IIIIW%
,/111 L V
.01
.'"
-- "
AI3F e dend- +
AI-AI,Fe eat
AI den~. +" - - ~ - J / ~ " "-s "-A/F~ d e n d ~ 1 2 1 ; e ;ut" ~l-Al,,Feeut,l ~ t ~ 3 -, " 3 " ". / , ~ P ~ ' I . AI-AI,'}FeeutI,,, ,
.001 0
1
2
3
" 4
Fe-content [at%] Figure 3.12. Comparison between experimental (full lines) and predicted (broken lines) solidification microstmcture selection map for Al-rich A1-Fe alloys (Gilgien et al. 1995).
H. Jones
36
lOO
104
10 ~
106
10"
"i" K/s Figure 3.13. Grain size vs. cooling rate during solidification of [[5]] Ti-6wt%Al4wt%V, (A) Ni-5wt%A1, (~) Fe-6.3wt%Si, and ( 9 type 316 stainless steel (Jones 1991). Full line represents a prediction (Boswell and Chadwick 1977) for aluminium.
This is illustrated in Fig. 3.18 (Kim et al. 1990) for an Al-based A1-Y-Ni-M (M = Mn or Fe) metallic glass. Metallic glasses exhibit a well defined fatigue limit corresponding to ,~0.35 of the flow stress. Fatigue life above this limit tends to be shorter than for crystalline materials because absence of work hardening allows cracks to initiate at relatively low stresses. Their wear resistance can be superior to that of crystalline materials of the same hardness under sliding conditions, but tends to be inferior under abrasive wear conditions. Metallic glasses undergo considerable creep deformation at temperatures close to but below their glass transition temperature Tg compared with crystalline metals and ceramics, Newtonian viscous flow is observed at much lower temperatures and up to higher stresses. Correlations between minimum creep rate and rupture time and between applied stress, minimum creep rate and temperature are shown in Fig. 3.19 (Rama Rao and Radhakrishnan 1984, Gibeling and Nix 1978) for an F e - N i - P - B glass. The homogeneous flow behavior of metallic glasses at temperatures just below Tg can be applied to hot work them into useful shapes. The increased scattering of electrons by the disordered structure and mixed population of atomic species in a metallic glass gives rise to high electrical resistivity, two or three times that of the same composition in crystalline form. In addition metallic glasses based on ferromagnetics such as iron, cobalt or nickel exhibit excellent soft ferromagnetic behavior. Nanocrystalline counterparts can exhibit superior soft or hard ferromagnetic
Rapid Solidification
37 0.5 w t . %
Dendritic
10 ~
Cellular 10-' 10 -2
~L
2wt. % 10 ~
,~
10"
t., 10 2 ~=~ e,~
10~
$=I
101
~ o~t ~g 02
4wt.
%
!
10"2
l . . . . . . . . . . . . . . . . . . ]
10 ~
..... ~ ~, . .
6wt.%
.
10"1
"~
10"2
"..................
1 8wt.
10 ~
%
10-1
10 .2 -
10 ~
10"1
,
10 .2
Velocity (cm/sec) Figure 3 . 1 4 . a - A 1 cell spacing vs. velocity of solidification front for five A1-Fe alloy compositions compared with predictions of tip stability (dashed lines) and array model (shaded regions) (Lu e t al. 1 9 9 4 ) .
behavior depending on their constitution and alloy compositions. Excellent corrosion resistance can also be developed by appropriate levels of alloying with chromium along with additions such as Mo, W, P or B. Catalytic effectiveness can also be enhanced for the glassy state compared with the crystalline condition. Crystalline materials made by rapid solidification can also show high hardness and strength by combinations of fine grain size, presence of fine disperoids/precipitation and increased solid solution hardening, without the reduced stiffness associated with the glassy state. The Hall-Petch plot shown in Fig. 3.20 (Nagahuma et al. 1993) for rapidly solidified A1-Ni-Mn(-Zr) alloys compared with wrought unalloyed aluminium illustrates both the increased direct contribution of particle and solid solution hardening ('~500 MPa) and the indirect effect of stabilizing small grain size with an increased Hall-Petch slope
38
H. J o n e s
104
I
i V
103
i
w
i
I
i
i
,
~
I
I
I
I
i
i
=
I
I
1,
i
~'.
~2 102 l.Lm 10 I
0.1 I
0.01
10"s
I
I
.
I
!
I
I
100
"1" K/s
10 s
t
101~
Figure 3.15. Dendrite cell size ~-2 vs. cooling rate J" during solidification for A1-4 to 5wt%Cu (open points) and A1-7 to 11 wt%Si (filled points). Full line represents ),2 j"l/3 = 50 # m (K/s) 1/3 (Jones 1984).
10
'"~1
o . Eote0t,0 ~m
x
Chadwick
196
9 Cooksey .tal
01
-
A
Elliott
+ L i v i n g s t o n et al
_
o
Clark&Elliott
0.011 0.001
0.01
r~__
. 1967
1968
13 B u r d e n & Jones 1970 -
9 Zimmermann I r ill
J
1964
Ibaraki&Okamoto M9ore&
~
9
1970 1976
et al 1988 I ~ I,f
0.1
,
,
,
,! 1
V mm/s
10
100
1000
Figure 3.16. Eutectic interphase spacing ,~ vs. V for the regular (lamellar) or a-A1A12Cu eutectic (Jones 1984, Zimmermann et al. 1989).
39
Rapid Solidification
..ii.iiii.i.!.ii!iiii!i!i!!!i!!ii!i!!!ii!;ii!i!!!i!!!i!i!iiii!i!i!i!!ii!i!i!i:!!i i~ii~iiii~i!i!ii!ii~i~ii~i~ii!~i!i!i~!ii~ii~i!i~i~i~i~i~ii~ii~ii~i ;i~!i~i~!ii!ii~i!~i~ii~i~ii~ii i .....................................................................................................................................
~!i!iii!i!iiiiiiiiiiiiiiiiiiii!iiiii!iiiii!iiiiiiiiiiiiiiii iiiiliiiiiiiii!iiiiiiiiiiiii!iiiiiiiiiiiiiiiiiiiiiiiiiii iiiiii!i!iiiiiiiiiiiiiiiii!iiiiiiiiiiiiiiiiiiiiii!i
%!~i (,,.~~ ,;
5:.i% 9,,,,,r
Figure 3.17. Shear band offsets on the outer surface of a bent metallic glass ribbon (Davis 1976).
6 0 0 _ AleeY 2 N i 9 M 1 "r" 4 0 0
~
1~
ol21 Mn
@11 Fe * brittle
200 EL (..9 uJ
60, 50 I
,
I
I
I
I
1400
d3
1200
EL
-
1000
.....
3~ 800
w
2
6
~b
2'o V~
3'o
4'0
(%)
Figure 3.18. Hardness Hv, Young's modulus E, fracture strength af and elongation to fracture Ef of A1-2at%Y-9at%Ni-lat%M (M = Mn or Fe) metallic glass ribbon vs. volume fraction Vf of fcc o~-A1 nanocrystals introduced into the metallic glass matrix (Kim et al. 1990).
40
H. Jones
10%
a)
~i~i I
I
I
"i
I
O
10-5 -
10-6-
eo
10-7 102
1
103
200
b)
I
-
c,, 1 0 0
-
I
10/+ t,ls)
m
E E
' "
105
,
106
9
b
50_.22
,l
'23
,I
,
I
-2z~ -25 P= tog[~s]- O03t+T
1
-26
Figure 3.19. (a) Minimum creep rate ks vs. rupture time tr for Fe-40at%Ni-14at%P6at%B metallic glass for stress range 600-1600 MPa (temperatures: ( 9 523 K, (O) 573 K); (b) applied stress o- vs. parameter P (ks is the strain rate in s -1 , T is the test temperature in K) for the same glass. From (Rama Rao and Radhakrishnan 1984) with data from (Gibeling and Nix 1978).
(3.7 MPa mm 1/2 compared with 2.3 MPa mm 1/2 for the wrought unalloyed condition). The new alloy chemistries afforded by rapid solidification have also led to improved thermal stability, such as in the A1-Fe-V-Si (8009) alloy developed by Allied Signal. Figure 3.21 (Gilman 1990) shows that 8009 retains relatively high strength in the range 150450~ where conventional wrought A1-Cu (2219) and A1-Zn-Mg-Cu (7075) lose most of their high ambient temperature strength as a result of precipitate coarsening. The re-
Rapid Solidification 1200
1000
(3.
8O0
:
O
1
41
i
AI-Ni-Mm-Zr
(~
AI-Ni-Mm
[-]
Fujila and Taba/a
/~
Wyrzykowski and Grabski
1
Hansen
-
t-"
600
-o
400
200
0
20
I
I
I
40
60
80
d 1/2
t
100
,,,
[
120
140
/ mm
Figure 3.20. Proof strength o-0.2 vs. inverse square root of grain size d for rapidly solidified A1-Ni-Mn(-Zr) alloys ((30) compared with results for wrought unalloyed aluminium ([]Am) (Nagahuma et al. 1993).
markable resistance to coarsening of the high volume fraction of the nanoscale silicide dispersoid in 8009 thus raises maximum service temperature by some 200 K. Creep resistance and corrosion resistance are also notably improved relative to established wrought ingot alloys of comparable room temperature strength. The limit on alloying set by incidence of unacceptable ingot macrosegregation in wrought ingot processing has been relaxed by rapid solidification processing for a number of other classes of material such as silicon irons for transformers, tool and bearing steels, nickel based superalloys and copper alloys for conductor applications. It is also a route of choice for processing the new class of high performance Nd-Fe-B permanent magnet materials based on the remarkable properties of the Nd2Fe14B compound.
3.4.2 Applications of rapid solidification It is worth recalling that modem interest in rapid solidification processing stems from curiosity-driven research in the 1960s aimed at establishing its capabilities for generating new alloy phases (Duwez 1967) and size-refined microstructures (Grant 1970). There is no doubt that it has contributed much to our present understanding of what governs formation of alloy phases, their mode of occurrence in solidification microstructures and their effect on properties and performance. It has also demonstrated the development potential of a number of previously underemployed production processes such as melt-atomization and spray forming, containerless processing, melt-spinning and thin-strip casting, and surface melt-traversing, all of which are undergoing intensive appraisal. A number of new
H. Jones
42
Tem~rature (*F) 0 o,. 5OO v
100
200
300
400
500
600
700
800
900
"
70
~,~oo
50
q,)
ta,,,. W
80O9
cO 3OO
,,.'K
c
co
40 N tll
3O
N. 200 2219-T851 ~ A
7075-T651
0
9
20
.
~o
9
0
~: 113 ~
O ~
9O
6O0
8O "~
/rr
soo
~' 4oo w
o~
t~
2219-T851
7075-T651
E lOO
o.,.., ...,.
!
80
r
"-
1
I
i
"
!
'
9
9
"
1
'"I
9
"
!
9
1
I
"'
' I
.
!
' " '"
.
I
"
i
I --
|
9
1_
"
I[
"
0 80
7O 7075-T651
-~ 6 o 9
j
C:~ " c ~0 -
c: ~ I-.
l
/
30 20
60
/
c
V)
J
5O
2219-T851
40 3O 8009
20 -
10
0
----
9 !
70 -
IJJ
CD
2O N E 10
0
so
c
o
~
.~
60
c ID 3O J--.
cll I..- 2OO
-
.~
5O m
8OO9
30O
70
50
10
100 150 200 250 300 350 400 450
Temperature (*C) Figure 3.21. Tensile properties of rapidly solidified 8009 (A1-Fe-V-Si) alloy vs. test temperature compared with high strength wrought alloys 2219-T851 and 7075-T651 (Gilman 1990).
R a p i d Solidification
43
materials have been developed. These include a range of new soft (Smith 1993) and hard (Guthrie 1993) magnetic materials for application in power distribution as well as in magnetic and electronic devices. New high performance light alloys have been established for both wear resistance (Amano et al. 1989) and high temperature (Gilman 1990) performance. Materials with enhanced catalytic performance (Hashimoto 1997) and for fuel cells (Kawashima et al. 1996) have been developed. Powder metallurgy tool steels and superalloys made from prealloyed atomized powder have an established market for inserts, engineering parts and hard facings (Abraham 1991). Direct spray deposition has been employed to produce mill rolls with two or three times the life of conventionally produced rolls (Ikawa et al. 1990). Rapid solidification has been used to develop new alloys for medical implants (Wang et al. 1988) and to make improved Ag-based feedstock particulate for producing dental amalgams (Vero et al. 1992). In addition rapid solidification has provided feedstocks for other developmental technologies such as mechanical alloying, plasma and high velocity spraying, and manufacture of metal matrix composites. It has also provided cost-effective routes to make steel fibres for reinforcement of concrete and castable refractories (Edgington 1977) and aluminium flake for imparting conductivity to plastics (Holbrook 1986) as well as to form ductile brazing foils for a wide variety of precision joining applications (Rabinkin and Liebermann 1993). The newest application, that of a metallic glass golf club head (Ashley 1998) with strength and hardness twice that of cast stainless steel or titanium but of lower modulus and intermediate density, made by bulk undercooling prior to solidification, would have been inconceivable 30 years ago.
REFERENCES
Abraham, T. (1991) Internat. J. Powder Metall., 27, 161. Alonso, J. A., Gallego, L. J. and Lopez, J. M. (1988) Phil. Mag., A58, 79. Amano, N., Odani, Y., Takeda, Y. and Akechi, K. (1989) Metal Powder Rept., 44, 186. Anderson, I. E. (1991) Adv. Mater. Process., 140(1), 30. Anon (1989) Adv. Mater. Process., 135(6), 12. Ashley, S. (1998) Mech. Eng., 120(6), 72. Boettinger, W. J. (1982) in Rapidly Quenched Metals, Vol. I, eds. Masumoto, T. and Suzuki, K. (The Japan Institute of Metals, Sendai, Japan), p. 85. Boettinger, W. J., Shechtman, D., Schaefer, R. J. and Biancaniello, E S. (1984) Metall. Trans., A15, 55. Boettinger, W. J., Bendersky,L. A., Coriell, S. R., Schaefer, R. J. and Biancaniello, E S. (1987) J. Cryst. Growth, 80, 17. Boswell, E G. and Chadwick, G. A. (1977) Scripta Metall., 11, 459. Boulby, K. A. and Wood, J. V. (1986a) Powder Metall., 29, 33. Boulby, K. A. and Wood, J. V. (1986b) Metal Powder Rept., 41, 299. Davies, H. A., Aucote, J. and Hull, J. B. (1974) Scripta Metall., 8, 1179. Davis, L. A. (1976) in Rapidly Quenched Metals, eds. Grant, N. J. and Giessen, B. C. (MIT Press, Cambridge, MA), p. 369. Duwez, P. (1967) Trans. Amer. Soc. Metals, 60, 607. Edgington, J. (1977) in Fibre-Reinforced Materials (Institute of Civil Engineers, London), p. 129. Feurer, U. and Wunderlin, R. (1977) DGM Fachbereicht, No. 38. Gelinas, C., Angers, R. and Pelletier, S. (1988) Mater. Lett., 6, 359. Gibeling, J. C. and Nix, W. D. (1978) Scripta Metall., 12, 919. Gilgien, E, Zryd, A. and Kurz, W. (1995) Acta Metall. Mater., 43, 3477. Gilman, E (1990) Metals and Materials, 6, 504. Grant, N. J. (1970) Fizika, 2 Suppl. 2, 16.1.
44
H. J o n e s
Greer, A. L. (1994) Nature, 368, 688. Guthrie, M. S. (1993) in Rapidly Solidified Alloys: Processes, Structures, Properties, Applications, ed. Liebermann, H. H. (Marcel Dekker, New York), p. 665. Hashimoto, K. (1997) Mater. Sci. Eng. A, A226/228, 891. Herlach, D. M. (1994) Mater. Sci. Eng. Rept., R12, 177. Herlach, D. M., Gillesen, E, Volkmann, T., Wollgarten, M. and Urban, K. (1992) Phys. Rev. B, 46, 5203. Herlach, D. M., Cochrane, R. E, Egry, I., Fecht, H. J. and Greer, A. L. (1993) Internat. Mater. Rev., 38, 273. Holbrook, A. L. (1986) Internat. J. Powder Metall., 22, 39. Hunt, J. D. (1979) in Solidification and Casting of Metals (The Metals Society, London), p. 3. Hunt, J. D. and Lu, S.-Z. (1996) Metall. Trans. A, 27A, 611. Ikawa, Y., Itami, T., Kumagai, K., Kawashima, Y., Leatham, A. G., Coombs, J. S. and Brooks, R. G. (1990) ISIJ Internat., 30, 756. Inoue, A., Zhang, T. and Masumoto, T. (1993) J. Non-Cryst. Solids, 156-158, 473. Jackson, K. A. and Hunt, J. D. (1966) Trans. Metall. Soc. AIME, 236, 1129. Johnson, W. L. (1996) Curr. Opin. Solid State & Mater. Sci., 1, 383. Johnson, W., Mamalis, A. G. and Hunt, H. (1976) Metallurgia and Metal Forming, 43(3), 68. Jones, H. (1984) J. Mater. Sci., 19, 1043. Jones, H. (1988) Mater. Lett., 6, 181. Jones, H. (1991) in Proceedings of the Second Japan International SAMPE Symposium, eds. Kimpara, I., Kageyama, K. and Kagawa, Y., SAMPE, p. 306. Juarez-Islas, J. A. and Jones, H. (1987) Acta Metall., 35, 499. Kawashima, A., Habazaki, H., Asami, K. and Hashimoto, K. (1996) Sci. Rept. RITU A, A42, 91. Kim, W. T. and Cantor, B. (1994) Acta Metall. Mater., 42, 3045. Kim, Y. H., Inoue, A. and Masumoto, T. (1990) Mater. Trans. JIM, 31,747. Kui, H. W., Greer, A. L. and Tumbull, D. (1984) Appl. Phys. Lett., 45, 615. Kurz, W. and Fisher, D. J. (1981) Acta Metall., 29, 11. Kurz, W., Giovanola, B. and Trivedi, R. (1986) Acta Metall., 34, 823. Kurz, W., Giovanola, B. and Trivedi, R. (1988) J. Cryst. Growth, 91, 123. Lu, S.-Z., Hunt, J. D., Gilgien, P. and Kurz, W. (1994) Acta Metall. Mater., 42, 1653. Lux, B. and Hiller, W. (1972) Praktische Metallographie, 8, 218. Maringer, R. E. and Mobley, C. E. (1974) J. Vac. Sci. Technol., 11, 1067. Massalski, T. B., Kim, Y. W., Vassamillet, L. E and Hopper, R. W. (1981) Mater. Sci. Eng., 47, 1. Midson, S. E and Jones, H. (1982) in Rapidly Quenched Metals, Vol. II, eds. Masumoto, T. and Suzuki, K. (The Japan Institute of Metals, Sendai, Japan), p. 1539. Moir, S. A. and Jones, H. (1991) Mater. Lett., 12, 141. Nagahuma, H., Ohtera, K., Higashi, K., Inoue, A. and Masumoto, T. (1993) Phil. Mag. Lett., 67, 225. Oguchi, M., Inoue, A., Yamaguchi, H. and Masumoto, T. (1990) Mater. Trans. JIM, 31, 1005. Ohnaka, I. (1985) Internat. J. Rapid Solidification, 1, 219. Overshott, K. J. (1979) Electron. Power, 25, 347. Peker, A. and Johnson, W. L. (1993) Appl. Phys. Lett., 63, 2342. Pelletier, S., Gelinas, C. and Angers, R. (1990) lnternat. J. Powder Metall., 26, 51. Perepezko, J. H. (1984) Mater. Sci. Eng., 65, 125. Prasad, Y. V. S. S., Ramachandrarao, E and Chattopadhyay, K. (1984) Acta Metall., 32, 1825. Rabinkin, A. and Liebermann, H. H. (1993) in Rapidly Solidified Alloys: Processes, Structures, Properties, Applications, ed. Liebermann, H. H. (Marcel Dekker, New York), p. 691. Raman, R. V., Patel, A. N. and Carbonara, R. S. (1982) Prog. Powder Metall., 38, 99. Rama Rao, E and Radhakrishnan, V. M. (1984) in Metallic Glasses, ed. Anantharaman, T. R. (TransTech, Aedermannsdorf, Switzerland), p. 225. Schmitt, H. (1979) Powder Met. Internat., 11, 17. Smith, C. H. (1993) in Rapidly Solidified Alloys: Processes, Structures, Properties, Applications, ed. Liebermann, H. H. (Marcel Dekker, New York), p. 617. Spaepen, E (1987) in Undercooled Alloy Phases, eds. Collings, E. W. and Koch, C. C. (TMS, Warrendale, PA), p. 187. Vero, B., Fauszt, A., Tardy, E, Takacs, J., Konczos, G. and Lovas, T. (1992) Mater. Sci. Technol., 8, 645. von Allmen, M., Huber, M., Blatter, A. and Affolter, K. (1984) Internat. J. Rapid Solidification, 1, 15. Wang, K. K., Gustavson, L. J. and Dumbleton, J. H. (1988) Mod. Dev. Powder Metall., 20, 361.
R a p i d Solidification
45
Weber, J. K. R., Krishnan, S. and Nordine, R C. (1991) J. Metals, 43(7), 8. Zimmermann, M., Carrard, M. and Kurz, W. (1989) Acta Metall., 32, 3305.
SELECTED BIBLIOGRAPHY
Ananthararman, T. R. (ed.) (1984) Metallic Glasses, Production, Properties and Applications (TransTech, Aedermannsdorf, Switzerland), 300 pp. Anantharaman, T. R. and Suryanarayana, C. (1987) Rapidly Solidified Metals (TransTech, Aedermannsdorf, Switzerland), 260 pp. Cahn, R. W. and Greer, A. L. (1996) in Physical Metallurgy, eds. Cahn, R.W. and Haasen, P., fourth edition (Elsevier Science, Oxford), p. 1723. Gilman, J. J. and Leamy, H. J. (1978) Metallic Glasses (American Society for Metals, Metals Park, OH), 350 PP. G~intherodt, H. J. and Beck, H. (1981) Glassy Metals I (Springer-Verlag, Berlin), 270 pp. Gtintherodt, H. J. and Beck, H. (1983) Glassy Metals II (Springer, Berlin, Germany), 400 pp. Herman, H. (ed.) (1981) Ultrarapid Quenching of Metals and Alloys (Academic Press, New York), 448 pp. Jacobson, L. A. and McKittnick, J. (1994) Mater. Sci. Eng. R, R l l , 355. Jones, H. (1982) Rapid Solidification of Metals and Alloys (Institution of Metallurgists, London), 83 pp. Jones, H. (1984) J. Mater. Sci., 19, 1043. Jones, H. (1991) Metals and Materials, 7, 486. Liebermann, H. H. (ed.) (1993) Rapidly Solidified Alloys: Processes, Structures, Properties, Applications (Marcel Dekker, New York), 808 pp. Luborsky, E E. (ed.) (1983) Amorphous Metallic Alloys (Butterworths, Stoneham, MA), 548 pp. Otooni, M. A. (ed.) (1997) Elements of Rapid Solidification (Springer-Verlag, Berlin), 320 pp. Sahm, P. R., Jones, H. and Adam, C. M. (eds.) (1986) Science and Technology ofthe Undercooled Melt (Martinus Nijhoff, Dordrecht, The Netherlands), 304 pp. Srivatsan, T. S. and Sudarshan, T. S. (1993) Rapid Solidification Technology: An Engineering Guide (Technomatic., Lancaster, PA), 737 pp.
These volumes provide extensive reviews of their subject areas that were comprehensive at the time of their compilation. More specialized review articles also continue to appear but are too numerous to list here. In addition, the series of international conferences that began in Brela, Yugoslavia, in 1970 and reached its ninth at Bratislava in 1996 and the tenth in Bangalore, India in 1999 contain numerous useful review articles within their Proceedings (for the ninth conference see Mater. Sci. Eng. A, 15th June 1997 issue, Vols. A 226-228, 1110 pp). The proceedings of two recent NATO workshops are also useful: Undercooled Metallic Melts: Properties, Solidification and Metastable Phases, eds. D. M Herlach et al., Mater. Sci. Eng. A, 1994, A178(1,2), 320 pp, and Science and Technology of Rapid Solidification and Processing, ed. M. A Otooni, Kluwer, Dordrecht, 1995, 390 pp). International Journal of Rapid Solidification, Vol. 1, No. 1 (1984) to Vol. 9, No. 4 (1996), is the only journal dedicated to publication of research papers in the field. It continues with its new name International Journal of Nonequilibrium Processing with Vol. 10, No. 1 in 1997.
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Chapter 4 Mechanical Alloying 4.1.
4.2. 4.3.
Introduction Nomenclature The Process of Mechanical Alloying 4.3.1 Raw Materials 4.3.2 Process Control Agents 4.3.3 Type of Mills
4.3.3.1 SPEX Shaker Mills 4.3.3.2 The Planetary Ball-Mill 4.3.3.3 Attritor Mills 4.3.3.4 Commercial Mills 4.3.3.5 New Designs 4.4. Mechanism of Alloying 4.5. Consolidation 4.6. Synthesis of Non-Equilibrium Phases 4.6.1 Solid Solubility Extensions 4.6.2 Synthesis of Intermetallics 4.6.3 Disordering of Ordered Intermetallics 4.6.4 Nanocrystalline Materials 4.6.5 Solid-State Amorphization 4.6.6 Displacement Reactions 4.7. Powder Contamination 4.8. Modeling 4.9. Industrial Applications 4.9.1 Nickel-Base Alloys 4.9.2 Iron-Base Alloys 4.9.3 Aluminum-Base Alloys 4.9.4 Magnesium-Base Alloys 4.10. Concluding Remarks References
49 5O 51 51 52 52 53 53 53 53 54 55 56 58 59 62 64 66 67 71 74 75 76 77 79 79 80 80 81
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Chapter 4 Mechanical Alloying C. SURYANARAYANA
4.1. I N T R O D U C T I O N
Mechanical alloying is normally a dry, high-energy ball-milling technique and has been employed to produce a variety of commercially useful and scientifically interesting materials. This technique was developed around 1966 by John S. Benjamin and his colleagues at the Paul D. Merica Research Laboratory of the International Nickel Company (INCO). The main goal was essentially to combine the advantages of precipitation-hardening and oxide-dispersion strengthening in a nickel-based superalloy intended for gas turbine applications. Benjamin (1976, 1990)reviewed the historic origins of the process. This process was referred to as "milling/mixing", but a patent attorney for the INCO coined the term m e c h a n i c a l alloying.
The formation of an amorphous phase by mechanical grinding of an Y-Co intermetallic compound in 1981 (Yermakov et al. 1981) and in the Ni-Nb system by ball-milling of blended elemental powder mixtures in 1983 (Koch et al. 1983) brought about the recognition that mechanical alloying is a potential non-equilibrium processing technique. Beginning from the mid-1980s, a number of investigations have been carried out to synthesize a variety of stable and metastable phases including supersaturated solid solutions, crystalline and quasicrystalline intermediate phases, and amorphous alloys (Suryanarayana 1995a). Efforts were also under way since the early 1990s to understand the process fundamentals of mechanical alloying (Maurice and Courtney 1990). Additionally, it has been recognized that this technique can be used to induce chemical (displacement) reactions in powder mixtures at room temperature or at least at much lower temperatures than normally required (Heinicke 1984, McCormick 1995). Because of all these special attributes, this simple, but effective, processing technique has been applied to metals, ceramics, polymers, and composite materials. The important attributes of mechanical alloying are listed in Table 4.1. The different facets of this technology have been reviewed periodically (Gilman and Benjamin 1983, Weeber and Bakker 1988, Koch 1989, Koch 1991, Herr 1995, Suryanarayana 1996, Lu and Lai 1998, Murty and Ranganathan 1998). The literature available till 1994 has been collected together in an annotated Bibliography (Suryanarayana 1995a). Additionally, a number of conferences are devoted to the science and technology of mechanically alloyed materials (Arzt and Schultz 1989, Yavari and Desr6 1990, Froes and de Barbadillo 1990, de Barbadillo et al. 1993). Furthermore, proceedings of the annual International Symposia on 'Metastable, Mechanically 49
50
C. Suryanarayana
Table 4.1. Attributes of mechanical alloying Production of a fine dispersion of second-phase particles Extension of solid solubility limits Refinement of grain sizes down to the nanometer range Synthesis of novel crystalline and quasicrystalline phases Development of amorphous (glassy)phases Disordering of ordered intermetallics Possibility of alloying of difficult to alloy elements Inducement of chemical (displacement) reactions at low temperatures Scaleable process
Alloyed and Nanocrystalline Materials' (ISMANAM) are regularly published in Materials Science Forum published by TransTech Publications, Switzerland (Shingu 1992, Yavari 1995, Schulz 1996, Fiorani and Magini 1997).
4.2. N O M E N C L A T U R E
Two different terms are commonly used in the literature to denote the processing of powder particles in high-energy ball-mills. Mechanical Alloying (MA) describes the process when mixtures of powders (of different metals or alloys/compounds) are milled together. Material transfer is involved in this process to obtain a homogeneous alloy. On the other hand, milling of uniform (often stoichiometric) composition powders, such as pure metals, intermetallics, or prealloyed powders, where material transfer is not required for homogenization, has been termed Mechanical Milling (MM). Some investigators have referred to MM as Mechanical Grinding (MG). Since "grinding" is normally thought of as an abrasive machining process that involves mainly shear stresses and chip formation, the term "milling" is preferred to include the more complex triaxial, perhaps partly hydrostatic, stress states that can occur during ball-milling of powders (Koch 1991). It should be realized that MA is a generic term, and some investigators use this term to include mechanical milling/grinding as well. However, we will distinguish between these two terms by using MA or MM, depending on whether material transfer is involved or not during processing. Reaction-milling (RM) is the process developed by Jangg et al. (1975). In this process the powder is milled without the aid of any process control agent (see later for its function during milling) to produce fine dispersions of oxides and carbides in aluminum (Jangg 1989). This proprietary process of Eckart-Werke in Germany achieves the dispersion of carbides by adding lamp-black or graphite during milling of aluminum. Adjusting the oxygen content via close control of the milling atmosphere (oxygen, argon, nitrogen, air, etc.) produces the oxides. Thus, the final product of milling contains the aluminum phase along with A14C3 and A1203 and these alloys are given the trade name DISPAL. Even without intentional additions of graphite, fine dispersions of A14C3 and A1203 could be produced when a process control agent is added during milling of metal powders. Thus, Singer et al. (1980) observed the presence of 10-nm size dispersoids of A14C3 and A1203
Mechanical Alloying
51
in a hot-extruded sample after milling the aluminum powder in the presence of Nopcowax 22-DSP. Milling of metal powders in the presence of reactive solids/liquids/gases (enabling a chemical reaction to take place) is now regularly employed to synthesize metal oxides, nitrides, and carbides (Avvakumov 1986, Calka et al. 1996). Thus, milling of titanium in a nitrogen atmosphere has produced titanium nitride (Calka 1991, Suryanarayana 1995b); several other compounds have also been produced in a similar way. Milling of tungsten with carbon (graphite) has produced tungsten carbide (Radlinski and Calka 1991). Milling of metal powders with boron has produced borides, e.g., TiB2 (Calka and Radlinski 1990). Another variation of milling adopted is cryomilling (Luton et al. 1989), in which the milling operation is carried out in the presence of liquid nitrogen. This process produces 2-10 nm sized aluminum oxy-nitride particles which strengthen the aluminum powder. 4.3. T H E P R O C E S S OF M E C H A N I C A L
ALLOYING
The actual process of MA starts with mixing of the powders in the right proportion and loading the powder into the mill along with the grinding medium (generally steel balls). This mix is then milled for the desired length of time until a steady state is reached when the composition of every powder particle is the same as the proportion of the elements in the starting powder mix. Sometimes the powder is milled to an intermediate state either to form metastable phases or to achieve certain desired properties. (During MM, however, the objective would be to produce certain desired properties in the material, e.g., formation of nanostructured or amorphous phases, or disordering of ordered intermetallics since no alloying/homogenization is required to occur). The milled powder is then consolidated into a bulk shape and heat treated to obtain the desired microstructure and properties. Let us now look at the individual components mentioned above. 4.3.1 Raw materials
The raw materials used for MA are widely available commercially pure powders that have particle sizes in the range of 1-200/zm. But the powder particle size is not very critical, except that it should be smaller than the grinding ball size. This is because the powder particle size decreases exponentially with time and reaches a small value only after a few minutes of milling. These powders fall into the broad categories of pure metals, master alloys, prealloyed powders and refractory compounds. The oxygen content of the commercially pure metal powders ranges from 0.05-0.2 wt.%. Therefore, it becomes important to choose high-purity powders, especially if one is interested in studying phase formation, since more often than not, the nature and amount of impurities in the system decide the nature of the final phase formed and the chemical constitution of the alloy. Dispersion strengthened materials usually contain additions of carbides, nitrides, and oxides. Oxides are the most common and these alloys are known as oxide-dispersion strengthened (ODS) materials. In the early days the powder charge for MA consisted of at least 15 vol% of a ductile compressibly deformable metal powder to act as a host or a binder. However, in recent years, mixtures of fully brittle materials have been milled successfully, resulting in alloy formation (Koch 1991). Thus, the requirement of having a ductile metal
52
C. Suryanarayana
powder during milling is no longer necessary. Accordingly, ductile-ductile, ductile-brittle, and brittle-brittle powder mixtures are milled now a days to produce novel alloys. Mixtures of solid powder particles and liquids have also been milled in recent times (Ivanov 1992).
4.3.2 Process control agents
A process control agent (PCA) is added to the powder mixture during milling, especially when the powder mix involves a substantial fraction of a ductile component. The PCAs are mostly organic compounds, which act as surface-active agents. The PCA adsorbs on the surface of the powder particles and minimizes cold welding between powder particles and thereby inhibits agglomeration. The surface-active agents adsorbed on particle surfaces interfere with cold welding and lower the surface tension of the solid material. Because the energy required for milling is equal to the product of new surface area generated times the surface tension, a reduction in surface tension results in the use of shorter milling times and/or finer powders. A wide range of PCAs has been used in practice. These are mostly organic compounds used at a level of about 1-4 wt.% of the total powder charge and include stearic acid, hexane, oxalic acid, methanol, ethanol, acetone, heptane, Nopcowax-22DSP, octane, toluene, etc. (Suryanarayana 1998). Majority of these compounds decompose during milling, interact with the powder and form compounds, and these get incorporated in the form of dispersoids inside the powder particles during milling. Thus hydrocarbons containing hydrogen and carbon, and carbohydrates containing hydrogen, carbon, and oxygen are likely to introduce carbon and/or oxygen into the powder particles, resulting in the formation of carbides and oxides which are uniformly dispersed in the matrix. These are not necessarily harmful to the alloy system since they can contribute to dispersion strengthening of the material. The hydrogen subsequently escapes as a gas or is absorbed into the metal lattice on heating or sintering. The presence of air in the milling container or milling of the powders at very low temperatures also has been shown to minimize welding, most probably due to the increased brittleness of the powder particles at such low temperatures. Metal powders (with an fcc structure) milled in a hydrogen atmosphere have been found to become brittle and not stick to themselves or the container. The choice of a PCA for milling depends on the nature of the powder being milled and the purity of the final product desired. The nature and amount of PCA used during milling determine the final powder particle size and powder yield. It should also be realized that there is no universal PCA. One has to decide on a PCA by looking at the possible interactions between the metal and the components in the PCA.
4.3.3 Type of mills Different types of milling equipment are used to produce mechanically alloyed powders. They differ in their capacity, efficiency of milling and additional arrangements for cooling, heating, etc.
Mechanical Alloying
53
4.3.3.1 SPEX shaker mills Shaker mills such as SPEX mills, which mill about 10 g of the powder at a time, are most commonly used to investigate the fundamental aspects of MA and for alloy screening purposes. In this mill, the vial containing the powder to be milled and the grinding balls, is secured in the clamp and swung energetically back and forth several thousand times a minute. The back-and-forth shaking motion is combined with lateral movements of the ends of the vial, so that the vial appears to be describing a figure-8 or infinity sign as it moves. With each swing of the vial the balls impact against the sample and the end of the vial, both milling and mixing the sample. Because of the amplitude (about 5 cm) and speed (about 1,200 rpm) of the clamp motion, the ball velocities are high (on the order of 5 m/s) and consequently the force of the bali's impact is unusually great. Therefore, these mills can be considered as high-energy variety. 4.3.3.2 The planetary ball-mill Another popular mill for conducting MA/MM experiments is the planetary ball-mill (referred to as Pulverisette) in which a few hundred grams of the powder can be milled at a time. The planetary ball-mill owes its name to the planet-like movement of its vials. These are arranged on a rotating support disk and a special drive mechanism causes them to rotate around their own axes. The centrifugal force produced by the vials rotating around their own axes and that produced by the rotating support disk both act on the vial contents, consisting of material to be ground and the grinding balls. Since the vials and the supporting disk rotate in opposite directions, the centrifugal forces alternately act in like and opposite directions. This causes the grinding balls to run down the inside wall of the vialmthe friction effect--followed by the material being ground and grinding balls lifting off and traveling freely through the inner chamber of the vial and colliding against the opposing inside wall--the impact effect. Even though the linear velocity of the balls in this type of mill is higher than that in the SPEX mills, the frequency of impacts is much more in the SPEX mills. Hence, in comparison to SPEX mills, Fritsch Pulverisette can be considered lower energy mills. 4.3.3.3 Attritor mills Attritors are the mills in which large quantities of powder (from a few pounds to 100 lb) can be milled at a time. The velocity of the grinding medium is much lower (about 0.5 m/s) than in Fritsch or SPEX mills and consequently the energy of the attritors is low. Attritors of different sizes and capacities are available. The operation of an attritor is simple. The powder to be milled is placed in a stationary tank with the grinding media. This mixture is then agitated by a shaft with arms, rotating at a high speed of about 250 rpm (Fig. 4.1). This causes the media to exert both shearing and impact forces on the material. Although much less efficient than either the SPEX or Fritsch mills, the laboratory attritor works up to 10 times faster than conventional jar mills. 4.3.3.4 Commercial mills Commercial mills for MA are much larger in size than the mills described above and can process several hundred pounds of powder material at a time. Mechanical alloying for
C. Suryanarayana
54
Water-cooled stationary tank
Gas seal
Steel ball bearings Ball mill Figure ,1.1.
Rotating impeller
Arrangement of rotating arms on a shaft in the attrition ball-mill.
commercial production is carried out in ball-mills of up to 3,000 lb capacity. The milling time decreases with an increase in the energy of the mill. For example, a process that takes only a few minutes in the SPEX mill may take hours in an attritor and a few days in a commercial mill.
4.3.3.5 New designs Several new designs of mills have been developed in recent years for specialized purposes. These include the rod mills, vibrating frame mills, and the equipment available from Dymatron, Inc., Cincinnati, OH; Super Misuni NEV-MA-8 from Nisshin Giken, Co., Ltd., Tokyo, Japan (with the ability to control the temperature of milling from very low temperatures by spraying liquid nitrogen and up to a high temperature of 300~ by electrical heating); Uni-Ball-Mill from Australian Scientific Instruments, Canberra, Australia (it is possible to control the nature and magnitude of impact of the balls in this machine with the help of magnets); HE Mill from M.B.N. srl, Rome, Italy; and Zoz Maschinenbau GmbH,
Mechanical Alloying
55
Kreuztal, Germany. Some special equipment is also designed for specific laboratory applications. The operating details and special features of these mills can be found in the respective brochures available from the manufacturers. 4.4. MECHANISM OF ALLOYING During high-energy milling the powder particles are repeatedly flattened, fractured and rewelded. Whenever two steel balls collide, some amount of powder is trapped in between them. Typically, around 1,000 particles with an aggregate weight of about 0.2 mg are trapped during each collision (Fig. 4.2). The force of the impact plastically deforms the powder particles, creates new surfaces, and enables the particles to weld together. This leads to an increase in particle size. Since in the early stages of milling, the particles are soft (if we are using either ductile-ductile or ductile-brittle material combination), their tendency to weld together and form large particles is high. A broad range of particle sizes develops, with some as large as three times bigger than the starting particles. The composite particles at this stage have a characteristic layered structure consisting of various combinations of the starting constituents. With continued deformation, the particles get work hardened and fracture by a fatigue failure mechanism and/or by the fragmentation of fragile flakes. Fragments generated by this mechanism may continue to reduce in size in the absence of strong agglomerating forces. At this stage, the tendency to fracture predominates over cold welding. Owing to the continued impact of grinding balls, the structure of the particles is steadily refined, but the particle size continues to be the same. Consequently, the interlayer spacing decreases and the number of layers in a particle increases. As mentioned earlier, the rate of refinement of the internal structure (particle size, crystallite size, lamellar spacing, etc.) is roughly logarithmic with processing time and therefore the size of the stating particles is relatively unimportant. In a few minutes to an hour, the lamellar spacing usually becomes small and the crystallite (or grain) size is refined to nanometer (1 n m - 10 -9 m or 10/k) dimensions (Fig. 4.3). The ease with which this can be achieved is one reason why MA has been extensively employed to produce nanocrystalline materials (Koch 1993, Suryanarayana 1995c, Koch 1997). After milling for a certain length of time, steady-state equilibrium is attained when a balance is achieved between the rate of welding, which tends to increase the average particle size, and the rate of fracturing, which tends to decrease the average composite particle size. At this stage each particle contains substantially all of the starting ingredients, in the proportion they were mixed together and the particles reach saturation hardness due to the accumulation of strain energy. The particle size distribution at this stage is narrow, because particles larger than average are reduced in size at the same rate that fragments smaller than average grow through agglomeration of smaller particles. The average particle size obtained at this stage depends on the relative ease with which agglomerates can be formed by cold welding, fatigue and fracture strength of composite particles, and resistance of particles to deformation. From the foregoing it is clear that during MA, heavy deformation is introduced into the particles. This is manifested by the presence of a variety of crystal defects such as dislocations, vacancies, stacking faults, and increased number of grain boundaries. The presence of this defect structure increases the diffusivity of solute elements into the matrix. Further-
C. Suryanarayana
56
Figure 4.2. Ball-powder-ball collision of powder mixture during MA (Gilman and Benjamin 1983).
more, the refined microstructural features decrease the diffusion distances. Additionally, the slight rise in temperature during milling further aids the diffusion behavior, and consequently, true alloying takes place amongst the constituent elements. While this alloying generally takes place nominally at room temperature, sometimes it may be necessary to heat treat the mechanically alloyed powder for alloying to be achieved. This is particularly true when formation of intermetallics is desired.
4.5. CONSOLIDATION
Widespread application of mechanically alloyed powders requires production of the powder in tonnage quantities and also efficient methods of consolidating the powders into bulk shapes. All the consolidation methods generally used in powder metallurgy processes can also be used for mechanically alloyed powders. However, since the powder particles in the mechanically alloyed condition are stronger and smaller (typically a few microns, even though the grain size may be only a few nanometers) in size than those used in conventional powder metallurgy operations, some special precautions need to be taken to minimize their activity and high level of interparticle friction. Conventional consolidation of powder to full density through processes such as hot extrusion and hot isostatic pressing normally requires use of high pressures and high tem-
Mechanical Alloying
57
E r N ......, t,,..
~
N
r ,,.=.=, t~
Decreasing ball-to-powder ratio
t.,.=
13.
Milling time, h Figure 4.3. Schematic showing refinement of particle and grain sizes with milling time. The rate of refinement increases and the minimum particle (grain) size reached is lower at higher ball-to-powder weight ratios.
peratures for extended periods of time. Unfortunately, this results in loss of the benefits achieved due to the metastable effects or nanostructures obtained by MA. Therefore, novel and innovative methods of consolidating the mechanically alloyed powders are required. (See also Chapter 13 on "Powder Consolidation" by Joanna R. Groza in this book.) Mechanically alloyed powders have been consolidated into useful bulk shapes by several processes. Since the mechanically alloyed powders have a high hardness, cold compaction is not an option. Furthermore, ODS alloys have been found not to densify during simple sintering. The most common method of consolidation is hot compaction followed by hot extrusion, or by direct hot extrusion at temperatures greater than half the melting point. This process could be used if one is not concerned with the loss of metastable effects, such as the crystallization of the amorphous phase in the powder. Some of the consolidation methods used for mechanically alloyed powders, both on laboratory and industrial scales, include hot isostatic pressing, powder rolling, Ceracon processing, plasma activated sintering, electrodischarge compaction, and explosive forming (shock or dynamic consolidation methods). Combination of these techniques is also sometimes used (He and Ma 1996). Some investigations were also carried out to compare the efficiencies of the different consolidation techniques in retaining the material in a nanostructured condition (Suryanarayana et al. 1997). Retention of nanostructures seems to be most efficient when consolidation is carried out at very high pressures and when the powder is exposed to these high pressures (and temperatures) for only a short period of time, e.g.,
58
C. Suryanarayana
Table 4.2 Important milestones in the development of mechanical alloying/milling 1966 1981 1983 1988 1989 1990
Development of ODS nickel-base superalloys Amorphization of intermetallics Amorphization of blended elemental powder mixtures Synthesis of nanostructures Occurrence of displacement reactions Observation of disordering of intermetallics
in shock consolidation methods. These operations are usually carried out at temperatures considerably lower than those used for conventional powders. For example, while conventional y-TiA1 powders are hot isostatically pressed at about 1,100~ (Beddoes et al. 1992), mechanically alloyed y-TiA1 powders could be consolidated to full density by hot isostatically pressing them at about 725~ (Suryanarayana et al. 1997).
4.6. S Y N T H E S I S OF N O N - E Q U I L I B R I U M P H A S E S
MA of suitable alloy compositions can result in the formation of a variety of equilibrium and non-equilibrium alloy phases. These include equilibrium and supersaturated solid solutions, stable and metastable crystalline and quasicrystalline intermediate phases, and glassy (amorphous) alloys (Koch 1991, Suryanarayana 1996). This technique is also capable of producing true alloys starting from pure elements that are not either easy to form by conventional means or sometimes even impossible to prepare, e.g., elements which are immiscible under equilibrium conditions. Rapid solidification processing (RSP) (rapid quenching of metallic melts at rates of ~106 K/s) also produces very similar microstructural and constitutional effects as MA (see Chapter 3 on "Rapid Solidification" by H. Jones in this book) and, therefore, comparisons are frequently made of the product phases obtained by these two techniques. However, MA processing is carried out entirely in the solid state. Thus, limitations imposed by phase diagrams, such as immiscibility in the liquid and solid states, do not apply to the MA process and therefore the alloying capabilities are likely to be much more extensive than by conventional ingot techniques or RSP methods. Furthermore, it should also be noted that whereas the technique of rapid solidification started as an academic curiosity in 1960 (Duwez et al. 1960, Duwez 1967) and progressed into an industrially mature technology about 20 years after the demonstration of its capabilities (Anantharaman and Suryanarayana 1987, Liebermann 1993), MA started as an industrial necessity in the late 1960s and only recently the scientific basis of MA is beginning to be understood. Some of the important milestones in the development of MMMM are listed in Table 4.2. Recent advances in understanding the "science" behind MMMM processing will be discussed first under the following headings: (1) Solid Solubility Extensions; (2) Synthesis of Intermetallics;
Mechanical Alloying (3) (4) (5) (6)
59
Disordering of Ordered Intermetallics; Nanostructures; Solid-State Amorphization; and Displacement Reactions.
The industrial applications of mechanically alloyed materials, such as oxidedispersion strengthened (ODS) alloys and others, will then be discussed towards the end of the Chapter.
4.6.1 Solid solubility extensions Solid solubility extensions have been achieved in many alloy systems by non-equilibrium processing methods such as RSP and vapor deposition. Similar results have also been reported in metal powder mixtures processed by MA. However, variation of solid solubility limits with the process parameters has not been systematically investigated. On mechanically alloying the blended elemental powder mixtures, interdiffusion between the two components occurs and solid solutions can form. This solid solubility is expected to increase with milling time as diffusion progresses and reach a saturation level, beyond which no further extension of solubility occurs. These can be considered as the solid solubility levels achieved by MA. Solid solubility levels have been generally determined from the changes in the lattice parameter values calculated from the shifts in peak positions in the X-ray diffraction patterns. There are certain difficulties associated with the solid solubility limits determined this way and, therefore, they may not be accurate for reasons discussed in an earlier article (Suryanarayana 1996). The absence of a second phase in the X-ray diffraction pattern is usually taken as an indication of formation of a homogeneous solid solution. But this may not always be true, since the appearance of the second phase in the X-ray diffraction pattern depends on the amount of the second phase, its particle size, the strain and defect content present in the material (this may shift the peak positions, sometimes asymmetrically). Therefore, caution must be exercised in interpreting the X-ray diffraction patterns from mechanically alloyed materials. Even though many of these factors have not been taken into account in many of the cases, solid solubility extensions, beyond the equilibrium values, have been reported in a number of alloy systems and some of these values are listed in Table 4.3. It should also be noted that in many cases equilibrium solid solutions (e.g., Cu-Ni, A1-Mg, Cu-A1) were synthesized by MA starting from blended elemental powders. Solid solutions (both equilibrium and metastable) form during MA because diffusion distances between particles are considerably reduced. The increased defect density and a slight rise in temperature during MA further aid diffusion. In fact, the temperature during MA seems to have a predominant effect. Increasing the temperature during MA tends to encourage decomposition of the solid solutions formed to reach thermodynamic equilibrium and oppose the effect of alloying that is possible by MA. The balance between these two opposing forces decides the final constitution of the alloy. For example, MA of a Cu-50 at. % Ag powder mixture at room temperature resulted in the formation of a homogeneous fcc solid solution. But MA at 473 K resulted in a fully decomposed two-phase
C. S u r y a n a r a y a n a
60
Table 4.3 Selected list of solid solubility extensions obtained by mechanical alloying Solvent
Solute
Maximum room temperature solid solubility (at.%) (equilibrium
by MA
(Massalski et al. 1990)) Ag A1
Cu
~0.0
100.0
Gd
~0.0
5.0
Cr
~0.0
Fe
Cd Co
Cu
Fe
0.025
5.0 4.5
Mg
1.2
23.0
Mn
0.3
18.5
Nb
--~0.0
15.0
Ru
~0.0
14.0
Ti
~0.0
Zn
3.5
3.0 ~50 6.0
C
~0.0
Cr
4.0
40
Cu
~0.0
100
Ag
~0.0
100
Hg
~0.0
70
Sb
1.2
3.7
Sn
~0.0
9.8
Ti
~0.0
AI
20.8
50
9.4
Cu
0.3
40
Mn
Co
4.0
50
Nb
AI
6.0
60
Ge
3.4
Ni
1.2
10
AI
7.2
27
C
"~0.0
12
Cr
22.5
50
Ga
10.8
50
In
~0.0
13
Nb
3.2
10
Ti
9.7
28
Ni
8.6
Ru
A1
0.7
14
Ti
A1
11.2
60
Cu
~0.0
10
Mg
~0.0
5O
Si
~0.0
8
W
Fe
2.0
40
Zr
A1
1.0
17.5
Ni
~0.0
7.0
Mechanical Alloying
61
mixture. Partially decomposed structures were obtained during milling at intermediate temperatures (Klassen et al. 1997). A similar conclusion was drawn by Ma et al. (1997) to explain the differences in the alloying behavior between two liquid immiscible systems Ag-Fe and Cu-Fe. Another interesting observation that has been made in mechanically alloyed materials is that the solid solubility achieved increases with increasing solute content in the starting blended elemental powder mixture, making it difficult to report a specific solid solubility value (Clark et al. 1995). This has been tentatively explained on the basis that the concentration gradients between the solvent and the solute are steeper at higher solute contents and that these are expected to result in increased diffusion and consequently in higher solubility levels. Neither the actual mechanism(s) for formation nor the limits for the solid solubility extensions in alloy systems obtained by MA have been well investigated. It was suggested (Schwarz et al. 1985) that the increased solid solubility of Ti in Ni in mechanically alloyed Ni-Ti powder mixtures was due to the metastable equilibrium between the ot-Ni solid solution and the Ni-Ti amorphous phase as opposed to the stable equilibrium between the ot-Ni and Ni3Ti phases. That is, the extent of supersaturation is limited to the composition where the amorphous phase starts forming. A similar conclusion was also arrived at by Murty et al. (1993). However, it has been noted in recent years that significant solid solubility extensions can be achieved in alloy systems (e.g., Ti-Mg) even when an amorphous phase did not form in them (Zhou et al. 1995). Thus, alternative explanations need to be sought to explain the formation of solid solutions and their limits of solid solubility. For example, it was suggested (Suryanarayana and Froes 1990) that formation of supersaturated solid solutions is related to the occurrence of nanocrystalline phases. The large volume fraction of grain boundaries in these materials is expected to enhance the diffusion and consequently extend the solid solubility levels in these types of systems. In support of this argument, it has been shown that the solid solubility of Mg in Ti is zero when the grain size of the titanium phase is on the order of microns, while significant solubility is reported when the titanium grain size is on the order of nanometers. Formation of (supersaturated) solid solutions by MA is more noteworthy in liquid immiscible systems. Accordingly, it has been reported that extensive solid solutions form in Cu-Co, Cu-Fe, Cu-Ta, and Cu-V systems. Gente et al. (1993) and Huang et al. (1994, 1997) suggested that the increased diffusivity due to the presence of structural defects and local stresses in nanocrystals is responsible for the formation of supersaturated solid solutions in these systems. Formation of homogeneous solid solutions is favored when the crystallite size of the constituents is reduced to below about 1-2 nm. The absolute limits of solid solubility extension in RSP alloys have been determined from thermodynamic considerations using the concept of To temperature at which, for a given composition, the solid and liquid phases have the same free energy. Supersaturated solid solutions can be obtained only when the liquid could be undercooled to a temperature below To (Boettinger and Perepezko 1993). Even though MA also is a nonequilibrium processing technique, there have not been any attempts so far to rationalize even the limited experimental data of solid solubility extensions; furthermore, the possible limits have not been defined theoretically. However, since no liquid phase is involved
62
C. S u r y a n a r a y a n a
during MA, it is doubtful if the same criterion as in RSP can be used to rationalize the solid solubility extensions achieved during MA. Even though the mechanisms by which supersaturated solid solutions form by MA and RSP techniques are different, comparisons have been frequently made, with no specific conclusions. For example, solid solutions have been formed in the whole composition range in the Cu-Ag system by both the techniques. A similar situation was obtained in several other systems by rapid solidification methods, but not by MA. Solid solutions have been obtained in the full composition range in the Cu-Fe, A1Sb-InSb, and Cu-Co systems by MA but not by RSP. The levels of solid solubility achieved are also different in different systems by the two different techniques (Suryanarayana 1999).
4. 6.2 Synthesis of intermetallics Because of the intimate mixing of the constituent metal powders during MA, it has been possible to synthesize a variety of intermetallic phases at appropriate compositions in a number of alloy systems. The intermetallics synthesized include both equilibrium and metastable crystalline phases and non-equilibrium quasicrystalline phases. Only a few reports exist on the synthesis and transformation behavior of quasicrystalline phases. Icosahedral phases have been synthesized by MA in Mg-Zn-A1 (Ivanov et al. 1989a), Mg-Cu-A1, A1-Cu-Mn (Eckert et al. 1989, 1990a), A1-Cu-MnGe (Bahadur et al. 1992), A1-Cu-Fe (Nasu et al. 1992, Asahi et al. 1994), A1-Cu-Ru (Politis et al. 1989), Ti-Ni-Fe-Si, and Mg-A1-Pd systems (Takeuchi et al. 1993, Suryanarayana 1996). These icosahedral phases have been found to be similar to those produced by RSP. It has been reported (Eckert et al. 1990b) that the nature of the phase formed was different depending on the milling intensity. At a very high milling intensity, a crystalline intermetallic phase formed while at a very low intensity, an amorphous phase formed. A quasicrystalline phase formed at intermediate intensities. The formation of the crystalline intermediate phase at the highest milling intensity was attributed to the increased temperature rise during MA resulting in the crystallization of the quasicrystalline phase. In support of this, it may be noted (Bahadur et al. 1996) that an as-milled Ti-Ni-Fe-Si amorphous alloy transformed to an icosahedral quasicrystalline phase on annealing at a temperature of 1025 K. This trend is somewhat contrary to what has been observed in RSP alloys. In RSP, a stable crystalline phase formed when the melt was cooled at low cooling rates and an amorphous phase formed at the highest cooling rates. A quasicrystalline phase formed at intermediate cooling rates (Suryanarayana and Jones 1988). A similar observation of formation of metastable phases at low milling intensities and stable phases at higher milling intensities has been reported in some other alloy systems also (Gerasimov et al., 1991, Guo et al. 1994, Suryanarayana et al. 1999). The equilibrium Hume-Rothery electron compounds--fl-, y- and E-brass phasesm were synthesized by MA of pure Cu and Zn powders mixed in the proper proportions (McDermott and Koch 1986, Martelli et al. 1988, Koch 1991). A number of intermetallics in other alloy systems have also been synthesized (Koch 1991, Suryanarayana 1995a). These include aluminides, silicides, and other intermetallics. Ceramic materials such as borides, carbides, nitrides, and oxides have also been synthesized by reaction milling. In
Mechanical Alloying
63
some cases, formation of the intermetallics seems to have taken place by a combustion synthesis-type reaction. Amongst these cases, there are also examples where the combustion took place only after "interrupted milling", i.e., milling of the powders for a given length of time, aging the powder at room temperature after stopping the milling, and then resuming the milling operation. An NiA1 compound was synthesized this way (Atzmon 1990). Similar results were also reported in other systems (Mukhopadhyay et al. 1996) and in fact these results compare very favorably with those obtained in the regular combustion synthesis process (Moore and Feng 1995). The synthesis of only a few metastable intermetallic phases has been reported by MA. A metastable deformation-induced martensite phase has been reported to form in mechanically alloyed Cu-Zn alloy powders (McDermott and Koch 1986). Similarly, synthesis of a bct-Fe phase on prolonged processing of iron powder in a nitrogen atmosphere has also been reported (Rawers et al. 1995). It is surprising that not many such deformationinduced metastable phases have been reported in mechanically alloyed powder mixtures in view of the fact that there are several instances where metastable phases are formed by cold working; a good example being the formation of deformation-induced metastable E (hcp) martensite in austenitic stainless steels. Recently, a metastable crystalline phase with rhombohedral crystal structure was reported to form in mechanically alloyed A1-Ge powders (Yvon and Schwarz 1993, Chattopadhyay et al. 1996). Formation of a metastable phase with a simple cubic structure was also reported in mechanically alloyed Te-Ag alloys (Chitralekha et al. 1995). Extreme caution should be exercised in identifying metastable crystalline phases produced by MA. It is possible that some of the phases formed are due to interstitial contamination, and these may be misinterpreted as new intermetallic phases. An accurate chemical analysis of the final resultant powder can, in many cases, clear up this confusion (Suryanarayana 1995b). Both disordered and ordered intermetallics (some of them metastable) have been synthesized by MA. In some cases, the intermetallics were synthesized directly by MA; but, in others an additional heat treatment was required after MA to form the intermetallic. For example, ordered intermetallics have been found to form directly on MA in Al-rich Al-transition metal systems. Some of the examples include A13Nb (Suryanarayana et al. 1994a), A15Fe2 (Mukhopadhyay et al. 1994), A13Zr (Suryanarayana et al. 1994b, Li et al. 1995a, Desch et al. 1991, 1996), (A1,X)3Zr, where X = Fe or Ni (Suryanarayana et al. 1994b, Li et al. 1995a), A13Hf, and (A1,X)3Hf where X = Fe or Ni (Li et al. 1995b). Reasons for the formation of ordered intermetallics have not been investigated in detail. It may be assumed that a phase will exist either in the ordered or disordered condition depending upon the balance between atomic disordering introduced by MA and the thermally activated reordering. The reordering is caused by the difference in energy between the ordered and disordered states. Thus, if this difference in energy is small, the alloy will exist in the disordered state whereas if it is large the alloy will be in the ordered state. It has been shown that both MA and MM produced the NiA1 phase in the ordered state while the FeA1 phase was found to be in the disordered state (Schropf et al. 1994). The ordering energy is related and scales up with the enthalpy of formation ( A H f ) and A H f values
64
C. Suryanarayana
for NiA1 and FeA1 are 72 and 25 kJ/mol, respectively (Schropf et al. 1994), confirming the above argument. A number of intermediate phases have been synthesized in the disordered condition. This is not surprising because MA involves heavy deformation and this destroys ordering in the lattice (Stoloff and Davies 1966). The number of disordered intermetallic phases synthesized is too numerous to list here and can be found in (Koch 1991, Suryanarayana 1995a, Lu and Lai 1998, Murty and Ranganathan 1998).
4.6.3 Disordering o f ordered intermetallics It has been long known that partially ordered phases are stronger than those wholly disordered or fully ordered (because at a certain value of the long-range order parameter, S, superdislocations separate into unlinked singles). Thus, it is of interest to study the mechanical behavior of materials in various states of partial order. Disordering phenomena of ordered alloys have also been studied to understand the mechanism of disordering and also to produce the disordered material that has a higher deformability than the ordered alloy. Alloys can be disordered by irradiation (Russel 1985), rapid solidification (Koch 1988), or heavy plastic deformation (Stoloff and Davies 1966). The present section will deal with the phenomenon of disordering of ordered intermetallics by MM. An exhaustive review on this aspect has been recently published (Bakker et al. 1995a). Mechanical milling introduces high energy into the material being processed. This energy can be stored in the material as atomic disorder and/or grain boundaries, i.e., A G (milling) = A G (disorder) + A G (grain boundaries). The atomic disorder in an intermetallic can be manifested in three different ways (Bakker et al. 1995a). First, the two atomic species involved can occupy the "wrong" sublattices and this is referred to as anti-site disorder. This type of disorder was observed in a number of mechanically milled compounds with the A15 structure, e.g., Ni3A1, Ni3Si, Fe3Ge and those with the B2 structure, e.g., CoGa. Second, triple defects can be generated. In CoGa, for example, vacancies on the Co-sublattice in combination with Co anti-site atoms in a ratio of 2:1 constitute the triple defects. Third, there could also be redistribution of interstitials wherein the interstitial atoms in the octahedral sites are transferred to the tetrahedral sites, e.g., Mn3Sn2, Fe3Ge2. Additionally, grain refinement increases the grain boundary area and this also raises the free energy of the system. The sum of the energy of these two effects (disordering and creation of grain boundaries) will be the total energy introduced into the material during milling. Mechanical milling of ordered intermetallics has been shown to result in one of the three following types of transformation (Bakker et al. 1995a, 1995b): (1) formation of a solid solution of one component in the other, i.e., the terminal solid solution based on the major component, e.g., Nb3A1, V3Ga, Ni3A1, Fe3Ge, Ni2V, and NbAu2; (2) formation of an amorphous phase, e.g., Nb3Sn, NiZr, NiV2, CoZr; and (3) formation of a different phase with a complex crystal structure, e.g., Ni3Sn2, TiSi2.
65
Mechanical Alloying CoZr
1.0
tv W IW
0.5CuTi
d c/ i3AI 0.0
M E C H A N I C A L A L L O Y I N G TIME
Figure 4.4. Schematic representation of variation of long-range order parameter S with milling time.
It has been noted that upon milling, the long-range order parameter (S) in the intermetallic is gradually reduced and, in many cases, the material may become totally disordered (S = 0), e.g., Ni3A1 (Jang and Koch 1990). In other cases, S is reduced with milling time but does not reach S = 0, i.e., partial order and partial disorder co-exist, e.g., CuTi (Seki and Johnson 1990). In other cases, the S value does not decrease at all and is maintained at S = 1; but, with continued milling the material becomes amorphous, e.g., CoZr (Cho and Koch 1993). These three situations are schematically represented in Fig. 4.4. Thus, upon milling, an ordered intermetallic can transform into a disordered crystalline phase (solid solution) or an amorphous phase either with or without complete loss of long-range order. If the product is a crystalline phase, the material has an extremely fine grain size, usually in the nanometer range. Disordering of intermetallics has been studied primarily by X-ray diffraction techniques. By measuring the intensity of the superlattice reflections relative to that of the fundamental reflections, the long-range order parameter S is evaluated. Additionally, it has also been noted that the lattice parameter increases slightly (0.3-0.8%) in the disordered state (Gialanella et al. 1992). This measurement of change in lattice parameter, of course, is not possible if there is a change in the crystal structure due to disordering. Mossbauer spectroscopy techniques, superconducting properties and magnetization methods have also been employed to study the disordering phenomena (Bakker et al. 1995a). Recently, Zhou and Bakker (1996) have demonstrated that magnetic measurements are very powerful in determining the nature of disordering if one of the atoms involved has magnetic moments.
66
C. Suryanarayana
During disordering of CoZr, Cho and Koch (1993) did not observe any decrease in the intensity of superlattice reflections (relative to fundamental lines) with milling time before amorphization occurred. It was thus concluded that the grain boundary energy in this system is high enough to drive the crystal-to-amorphous transition without any disordering occurring. Zhou and Bakker (1996) observed that magnetization in the mechanically milled Co-Zr alloys increased continuously with milling time up to about 40 h (no reduction in the relative intensity of superlattice reflections was observed) and also that the lattice parameter of the disordered nanocrystalline phase increased with milling time. From these two observations, they concluded that disordering of CoZr occurs by atomic disorder and grain refinement (i.e., creation of additional grain boundaries). Thus, a combination of techniques can provide confirmatory evidence and also an accurate account of the mechanism of disordering. Whether an intermetallic transforms to a solid solution or an amorphous phase on milling is determined by the relative free energy values of the amorphous and crystalline phases with respect to the energy stored in the intermetallic by mechanical milling. For example, the intermetallic will transform to the solid solution if the enthalpy of the amorphous state as estimated by the Miedema analysis (Bakker 1998) is higher than that of the solid solution. Thus, if AG (milling)> AG a-c, where AG a-c represents the difference in free energy between the ordered crystalline and amorphous phases, complete amorphization occurs. On the other hand, if AG (milling)< AG a-c, no amorphization occurs; instead, a solid solution would form. Similar conclusions can also be reached by observing the nature of the phase diagrams. If the phase diagram shows that the intermetallic transforms to a (disordered) solid solution before it melts, then milling of the intermetallic will result in the formation of a (disordered) solid solution. Continued milling may then produce the amorphous phase. But if the intermetallic melts congruently (irreversibly ordered or permanently ordered intermetallic), then milling of the intermetallic will produce an amorphous phase directly. This has been shown to be true in a number of instances (Suryanarayana 1995a, Bakker et al. 1995a).
4. 6.4 Nanocrystalline materials Nanocrystalline materials are single-phase or multiphase materials, the crystal size of which is of the order of a few (typically 1-100) nanometers. Because of the extremely small size of the grains, a large fraction of the atoms is located in the grain boundaries and thus the material exhibits enhanced combinations of physical, mechanical, and magnetic properties (compared to material with a more conventional grain size, i.e., > 1 #m). Nanocrystalline materials show increased strength, high hardness, extremely high diffusion rates, and consequently reduced sintering times for powder compaction. The reader is referred to some recent comprehensive reviews for full details of the processing, properties, and applications of these materials (Gleiter 1989, Siegel 1991, Koch 1993, Suryanarayana 1995c and Chapter 12, this volume). Nanocrystalline materials can be synthesized by a number of techniques starting from the vapor phase (e.g., inert gas-condensation), liquid phase (e.g., electrodeposition, rapid solidification), and solid state (e.g., mechanical attrition). The advantage of using MA
Mechanical Alloying
67
for synthesis of nanocrystalline materials lies in its ability to produce bulk quantities of material in the solid state using simple equipment and at room temperature. Plastic deformation generally occurs by slip and twinning at low and moderate strain rates, while at high strain rates it occurs by the formation of shear bands, consisting of dense networks of dislocations. Transmission electron microscopy investigations have shown that plastic deformation by ball-milling, as anticipated, occurs via shear bands which are approximately 0.5 # m wide (Hellstern et al. 1989). The plastic strain in the material increases due to increasing dislocation density in the early stages of ball-milling. At a threshold dislocation density, even at only moderately elevated temperatures, the material relaxes into subgrains separated by low-angle boundaries, leading to a decrease of atomic level strain. During subsequent milling the process of high deformation/subgrain formation is repeated resulting in the subgrains becoming finer and finer, and the relative orientation of the subgrains with respect to each other ultimately becoming completely random. Once the subgrains reach a critical level of refinement, further refinement becomes virtually impossible since the stresses required for dislocation movement are enormously high due to the Hall-Petch strengthening. Thus, nanostructures with a minimum grain size are produced. In the case of fcc metals, this minimum grain size was found to scale with the melting temperature, or bulk modulus, of the given element (Eckert et al. 1992). However, in the case of bcc and hcp metals the minimum grain size attained was essentially independent of the melting temperature of the metals (Koch 1993). Nanostructure processing leads to the formation of supersaturated solid solutions and intermetallic compounds, which are otherwise difficult to obtain. For example, a supersaturated solid solution can be obtained in a system, if the grain size is in the nanometer regime; this is irrespective of the method by which the material is produced in the nanocrystalline state. Accordingly, the solid solubility of Bi in Cu is increased from the equilibrium value of <10 -4 to 4 at.% in the nanocrystalline state (Birringer et al. 1988). Similarly, formation of intermetallics such as Pd3Bi and CuEr has been achieved by nanostructure processing. As mentioned above, the enhanced solid solubility limits in mechanically alloyed materials is mostly due to the increased volume fraction of atoms residing at the grain boundaries in nanostructured materials. Since all the aspects of nanostructures are covered in Chapter 12 on "Nanocrystalline Materials" in this book, this topic will not be further covered in this chapter. 4.6.5 Solid-state amorphization
Amorphous phases are formed in suitable alloy systems by processing techniques such as rapid solidification from the liquid state, vapor-deposition, and laser-processing (Cahn 1991). In these methods, there is a change in the state of matter, i.e., a solid phase is formed from either the liquid or the vapor phase, and the high effective "quenching" rate and associated undercooling have been shown to be responsible for the amorphization process. There are also methods of amorphizing a solid without passing through the liquid (or vapor) state at any stage. Known as solid-state amorphization reactions (SSAR), these include irradiation, interdiffusion of metallic elements, pressure-induced vitrification, and mechanical deformation. Two reports of amorphization have triggered an increased research activity in this area. These are the synthesis of an Ni-Nb amorphous
68
C. Suryanarayana
phase starting from blended elemental powders of Ni and Nb by MA (Koch et al. 1983) and amorphization of intermetallic compounds (e.g., YCo3, Y2Co7, YCo5) by MM (Yermakov et al. 1981, 1982). Since then, amorphous phases have been reported in a number of alloy systems by both MA and MM. Amorphous phases have also been obtained by MM of quasicrystalline phases produced by RSP (Zhang and Wang 1996). Amorphous phases have been synthesized by MA starting from either elemental (of two or more types) or by MM of intermetallic compound powders. A reasonably exhaustive listing of the alloy systems in which amorphous phases have been found to form by MA or MM can be found in Gaffet et al. (1993) and Suryanarayana (1999). The semimetals Si and Ge have also been partially amorphized by MM (Gaffet et al. 1993). A further point to be considered is whether, or not, there exists a difference in the way an amorphous phase forms by MA from a mixture of blended elemental powders (A and B) and by MM from the intermetallic compound (Am Bn). Two reaction routes are possible: (i) A + B --+ (AmBn)crystalline ~ (AmBn)amorphous (ii) (Am Bn)crystalline ~ (Am Bn)amorphous Instances of occurrence of amorphization by both the routes can be found in the literature (Koch 1991, Suryanarayana 1995a, 1999). In general, the end product is the same irrespective of whether the starting material is a mixture of the elements, a single intermetallic phase, or a mixture of two or more such phases. However, it has been reported that MgToZn30 can be amorphized only by MM of the intermetallic, and not by MA of the blended elemental powder mixture (Calka and Radlinski 1989). Amorphization in ordered alloys seems to follow the sequence (Jang and Koch 1990): Ordered phase ~ disordered phase (loss of long-range order) -+ fine-grained (nanocrystalline) phase ~ amorphous phase. As mentioned earlier, it appears that an intermetallic (either as the starting phase or that formed on MA of blended elemental powders) becomes amorphous when its free energy is raised above that of the amorphous phase. The number of alloy systems that show amorphization by MA or MM is too numerous to list here. Even though it appears that any alloy system can be made amorphous under appropriate conditions of milling, it may not be true. This is because a number of variables, the more important of them being milling energy, milling temperature, and impurity contamination, control the constitution of the final product. Increased milling energy (achieved by a higher ball-to-powder weight ratio, increased speed of rotation, etc.) is normally expected to introduce more strain and increase the defect concentration in the powder and thus lead to easier amorphization. However, higher milling energies also produce more heat (and higher temperatures) and this can result in the crystallization of the amorphous phase. Therefore, a balance between these two effects will determine the nature of the final product phase. Examples of the occurrence of both crystalline and amorphous phases at high milling intensities (energies) are available in the literature (Koch 1996).
M e c h a n i c a l Alloying
69
A few reports also exist which describe the crystallization of amorphous alloys by MA. Eckert et al. (1988) reported that if the milling intensity were too high, Zr-Ni alloys which were amorphized at a lower milling intensity would exhibit crystalline structures. This was explained as due to the temperature rise of the powder caused by the increased kinetic energy of the milling media. Similar results were also reported in other systems (Mizutani and Lee 1990, Trudeau et al. 1990, Trudeau 1994). Mechanical milling has also resulted in the crystallization of amorphous alloys in some cases. Crystallization of amorphous alloys by MA/MM, whether formed during milling or as the starting phase, has been referred to as mechanical crystallization; this behavior is not well understood. It has also been reported that a crystalline F-TiA1 phase formed on milling beyond the amorphous phase in Ti- A1 alloys (Oleszak et al. 1993). This has been, however, suggested to be due to interstitial contamination of the powder (Suryanarayana 1995b). There have also been conflicting results on the effect of milling temperature on the nature of the phase formed. Generally, a lower milling temperature was found to accelerate the amorphization process (Pathak et al. 1993, Yamada and Koch 1993). See Koch (1994) for a detailed review on the temperature effects during MA/MM. The mechanism of amorphization by MA and/or MM is not clearly understood. The early investigators (Yermakov et al. 1981, 1982) assumed that the powder particles melted because of the very high rate of plastic deformation, and consequent rise in the powder temperature. Subsequent rapid quenching of the liquid by heat conduction into the less deformed, and hence cooler, interior regions of the particles results in the formation of the amorphous phase (like in RSP). However, energy input calculations and temperature measurements suggest that the temperature rise is not large enough for the powder particles to melt. Researchers now believe that amorphization during MA is not purely a mechanical process and that a solid-state reaction similar to that observed in thin films (Schwarz and Johnson 1983) occurs during MA. During MM, destabilization of the crystalline phase is thought to occur by the accumulation of point and lattice defects, and anti-phase boundaries. Schwarz and Johnson (1983) proposed that formation of amorphous phases from two elements in the solid state requires that (i) the two elements have a large negative heat of mixing in the amorphous (liquid) state (alloy systems with a positive heat of mixing have also been amorphized by MA; but, negative heat of mixing is an essential condition for glass formation by RSP) and (ii) the two elements have very different atomic sizes, allowing one of the components to diffuse more rapidly than the other. The negative heat of mixing provides the thermodynamic driving force for the reaction and the large difference in the diffusivities is necessary for favoring the kinetics of the reaction. Thus, owing to these factors, the free energy of the amorphous phase is lower than that of the mixture of the pure crystalline compounds, thus explaining the formation of the amorphous phase by MA of a mixture of pure components. Beke et al. (1991a, b) estimated the elastic mismatch energy stored in an ordered solid solution when its long-range chemical order is destroyed. They showed that if Tc > Tm and the ratio of the elastic mismatch energy to the ordering energy is high enough, amorphization is expected to occur. This has been proved to be a valid criterion in many cases of amorphization observed by irradiation.
70
C. S u r y a n a r a y a n a
Recently, attempts have been made to predict the ability of an alloy to become amorphous under MA/MM conditions. Since both atomic sizes and heats of formation are important in the formation of amorphous phases, Zhang (1993) plotted (RA -- R s ) / R 8 (where RA and R8 are the atomic radii of the components A and B, respectively) against A H, heat of formation of the amorphous phase, and observed that the regions where an amorphous phase forms or does not form can be separated by a straight line given by the equation:
(RA -- RB = A R ) / R B
= 0.068 A H + 0.716.
The proportion of correct predictions for the glass-forming alloys was 89.2% while that for the non-glass-forming alloys was only 71.4%. A theoretical model was also proposed in which amorphization was assumed to be realized through interstitial impurity formation during MA (Chakk et al. 1994). It was assumed that amorphization occurred when impurity atoms penetrated into interstitial sites and distorted the lattice locally. When the local distortions reached some critical value, the long-range order of the lattice was destroyed (destabilization of the crystalline phase) and an amorphous phase formed. Even though amorphization has been observed by both MA and MM, there is no quantitative data available to define the "critical" energy input (or other parameters) which can be equated to the critical cooling rate which should be exceeded to form the amorphous phase in RSP studies. Modeling of MA/MM (currently in progress in different laboratories) can provide additional information in this direction. Significant differences have been observed in the formation of metastable intermediate phases or amorphous alloys by MA/MM and RSP techniques. Formation of an amorphous phase has been reported in the Ti-A1 system by MA (due to contamination?); not possible by RSP methods. An important difference noted between these two methods is that while amorphous phases form in the vicinity of eutectic compositions by rapid solidification, they form near the equiatomic composition by MA. Furthermore, the glass formation range (the compositions range over which glasses can form) is wider in the mechanically alloyed condition than in rapid solidification. Formation of nanostructures has been reported in many alloy systems by mechanical alloying methods, but not by rapid solidification methods. But, devitrification of amorphous phases obtained by rapid solidification has led to the synthesis of nanostructure composites useful for magnetic applications (Suryanarayana 1995c, Lu 1996). Another important difference between mechanically alloyed and rapidly solidified alloys containing dispersoids appears to be in the size and distribution of the dispersoids. Figure 4.5 shows a pair of transmission electron micrographs comparing the matrix grain sizes, and size and distribution of dispersoids, in HIP compacts of Ti3Al-based alloys containing Er203 dispersoids. A Ti3A1 alloy to which 2 wt.% Er was added was rapidly solidified and then HIPed at 850~ (Fig. 4.5a). The mechanically alloyed material was Ti-25Al-10Nb-3V- 1Mo (at.%) to which 2 wt.% Er had been added and the alloy powder was HIPed at 1,000~ (Fig. 4.5b). Even after HIPing at a higher temperature, in comparison to the rapidly solidified alloy, the mechanically alloyed material showed a finer matrix
Mechanical Alloying
71
grain size, more uniform distribution of the dispersoids, absence of large dispersoids at the grain boundaries, and absence of dispersoid-free zones near the grain boundaries. Majority of the above-mentioned and other differences between materials processed by these two techniques could be traced to the fact that rapid solidification is a liquid-tosolid transformation while MA is a fully solid-state transformation.
4. 6.6 Displacement reactions It was recognized in 1989 that MA may be used as the basis of a chemical refining process (Schaffer and McCormick 1989) by demonstrating that the pure metal Cu can be synthesized when CuO and Ca were ball-milled together at room temperature. Simultaneous reduction of CuO and ZnO by Ca has also been shown to result in the formation of flbrass. This method of metal synthesis directly from their oxides or chlorides has now been extended to the synthesis of a number of metals, e.g., Zn, Ti, Zr, Ta, Gd, Er, Sm, V, W and some intermetallics, e.g., SmCos. Results of the Australian group have recently been summarized in a review article (McCormick 1995). Other groups also have now taken up this type of work using MA. Most of the reactions studied to date have been displacement reactions of the type: MO + R ---~ M + R O where a metal oxide (MO) is reduced by a reductant (R) to the metal (M). Metal chlorides also have been reduced to metals this way. The displacement reactions studied by MA are characterized by a large negative free energy change at room temperature and are therefore thermodynamically feasible at room temperature. However, commercial operations by the pyrometallurgical techniques are conducted at elevated temperatures to overcome the kinetic barriers and achieve sufficiently high reaction rates. Mechanical alloying can provide the means to substantially increase the reaction kinetics of the displacement reactions. This is because during MA, the repeated fracturing and welding of powder particles increases the area of contact between the particles due to a reduction in particle size and also allows fresh surfaces to come into contact repeatedly. This allows the reaction to proceed without diffusion through the product layer. Additionally, the high defect densities induced by MA/MM accelerate the diffusion process. As a consequence, these reactions will now occur at room temperature. If a reaction cannot occur at room temperature, the particle refinement and consequent reduction in diffusion distances (due to microstructural refinement) can at least significantly reduce the reaction temperatures. The product of the displacement reactions normally consists of two phasesmthe metal (or a compound) and the oxide or chloride associated with the reductant. The removal of the unwanted reaction by-product can be difficult due to the high reactivity of the metal phase associated with nanocrystalline grain sizes and intermixing of the phases induced by the MA process. The by-product phase can be easily removed if the metal particles are embedded in a continuous matrix of the byproduct phase. Removal of the byproduct is achieved by leaching the product mixture in a dilute acid or hot water, or by vacuum
C. Suryanarayana
72
RS+HIP'ed
850~
,,/!~; ~~
" ,
[ .....
' MA + HIP'ed
1000~
B
b
Figure 4 5 Transmission electron micrographs showing the difference in the m a trix grain size, and size and distribution of dispersoids in rapidly solidified (RS) and mechanically alloyed (MA) materials
t
Mechanical Alloying
TIC14-
Kroll Process +Mg (!) @ 1000" C
73
Ti sponge Melt
_., %.
Ingot Casting PREP
Ti powder Figure 4.6. Schematic illustration showing the route by which titanium alloy powder can be synthesized directly from its tetrachloride by mechanical alloying with a reducing agent.
distillation. The use of carbon or hydrogen as the reductant produces gaseous C0 2 or water vapor as the byproduct and obviates the need for leaching/distillation. The displacement reactions have now been studied to synthesize metals, alloys, intermetallics, and rare-earth permanent magnet alloys. The rare-earth alloys synthesized include Nd2Fel4B, Sm2Fel7N2.6, Sm2Fel4Ga3C2, SmCo5, Sm2Co17, SmFeTi and SmCoFe. The product of MA in these cases is usually an amorphous phase that needs to be heat treated to crystallize the hard-magnetic phase. An important advantage of the MA process is that nanocrystalline microstructures are developed and these result in remanence-enhanced magnetic materials. Since with remanence enhancement there is no need for an alignment process, significant potential exists for achieving cost savings and simplification of the manufacturing process. The mechanically driven displacement reaction offers a number of advantages over conventional metal processing techniques. First, it enables the reduction of a number of oxides and halides to pure metals at room temperature, thus effecting energy savings. Second, if a number of components are reduced simultaneously, then it is possible to produce an alloy without first having to convert the oxides to pure metals and then to the desired alloy. Third, for powder metallurgy applications it allows the direct formation of powder product without first having to manufacture the bulk alloy and then convert it to powder form. Thus, a number of high-temperature processes can be combined into one single room-temperature process with the potential for significant cost savings (Fig. 4.6) (Schaffer and McCormick 1992).
74
C. S u r y a n a r a y a n a
4.7. P O W D E R C O N T A M I N A T I O N
A major concern in the MA process is the impurities that get into the powder and contaminate it. The small size of the powder particles, availability of large surface area, and formation of fresh surfaces during milling all contribute to the contamination of the powder. Thus, it appears as though powder contamination is an inherent drawback of the technique, unless special measures are taken to avoid/minimize it. As mentioned earlier, in the early stages of MA, the metal powder coats the surface of the grinding medium and the inner walls of the container. This was expected to have prevented contamination of the powder and so no attention was paid to the problem of powder contamination. However, when different results were reported by different groups of researchers on the same alloy system, it was recognized that contamination could be a serious problem. This problem appears to be ubiquitous and is now encountered in many investigations, especially when reactive metals such as titanium and zirconium are milled. The magnitude of contamination appears to depend on the time of milling, the intensity of milling, the atmosphere in which the powder is milled, and the difference in strength~ardness of the powder and the milling medium. Whereas, 1-4 wt.% Fe has been found to be normally present in most of the powders milled with the steel grinding medium, amounts as large as 7 wt.% (20 at.%) Fe in a W-C mixture milled for 310 h (Wang et al. 1997), 13 wt.% (33 at.%) Fe in pure W milled for 50 h in a SPEX mill (Yang et al. 1993), and 60 ac% Fe in a W-5 at.% Ni alloy milled for 60 h in a SPEX mill (Courtney and Wang 1992) were also reported. These are very high levels of contamination. Iron contamination has also been observed in other alloy systems. For example, the presence of 16 at.% Fe was reported in a Ta-30 at.% A1 alloy powder milled for 400 h in a ball-mill (E1-Eskandarany et al. 1991). Similarly, large amounts of oxygen (up to 36.5 at.%) and nitrogen (up to 22.6 at.%) have also been reported to be present in AI-Ti powders milled for 400 h in a low energy ball-mill (Saji et al. 1992). Contamination of metal powders can be traced to (i) chemical purity of the starting powders; (ii) milling atmosphere; (iii) milling equipment; and (iv) the process control agents added to the powders. Contamination from source (i) can be either substitutional or interstitial in nature, while contamination from source (ii) is essentially interstitial in nature and that from (iii) is mainly substitutional, even though carbon from the steel equipment can be an interstitial impurity. Contamination from the PCA leads to interstitial contamination. Several attempts have been made in recent years to minimize the powder contamination during MA. One way of minimizing the contamination from the grinding medium and container is to use the same material for the container and grinding medium as the powder being milled. In this case also there will be wear of the grinding medium and the debris gets incorporated into the powder. Thus, even though there is no contamination, the chemistry of the final powder will be quite different from the starting powder; the metallic content (from the container and balls) would be higher than in the initial powder. This can be compensated for if we know how much of the metallic content is increased. The above solution may be possible in some cases; but it is difficult in many cases due to the non-availability of the special grinding medium and containers.
Mechanical Alloying
75
The problem of milling atmosphere is serious and has been found to be the major cause of contamination in many cases. In fact, it has been observed that if the container is not properly sealed, the atmosphere surrounding the container, usually air (containing nitrogen and oxygen), leaks into the container and contaminates the powder. Thus, when reactive metals like titanium and zirconium are milled in improperly sealed containers, the powders are contaminated with nitrogen and oxygen. Attempts have been made to improve the container seal integrity to prevent the outside atmosphere leaking inside. Use of high-purity argon (99.998%) atmosphere and improvements of the seal quality resulted in the processing of high-quality titanium alloy powder with as little as 100 ppm oxygen and 15 ppm nitrogen (Goodwin and Ward-Close 1995). This process, however, may not be economically viable and hence may not be feasible on an industrial scale.
4.8. M O D E L I N G
From the description of the MA process it is easy to realize that MA is a process involving a number of both independent and interdependent variables. Like in any other process, modeling of MA is also carried out to identify the salient factors affecting the process and establish process control instrumentation. By modeling the process effectively, it is possible to bring down the number of actual experiments to be conducted to optimize the process and achieve a particular application. The number of variables involved in the MA process is very large. For a particular alloy system the variables include the type of mill, intensity of milling, type of milling media, ball-to-powder weight ratio, the atmosphere under which the powder is milled, purity of the powders, milling time, milling temperature, and nature and amount of the PCA used. All these have a significant effect on the constitution of the powder. Even on a local scale, the nature of impacts between two balls, the frequency of impacts, the amount of powder trapped between two balls during a collision could vary from point to point. Thus, modeling the MA process is a difficult task. In spite of this, some attempts have been made in recent times (Courtney 1994, Courtney and Maurice 1996, Magini and Iasonna 1995, Magini et al. 1996, Abdellaoui and Gaffet 1996) and moderate success has been achieved in modeling the mechanics of the process. From the actual experiments conducted, attempts have also been made to correlate the phases formed with the process parameters during milling. But, we are far from a situation where we can predict the final chemical constitution of the powder (type and description of phases formed). As Courtney and Maurice (1996) succinctly put it: "It is important to appreciate what can be expected of even the most successful efforts of this kind. 'Absolute' predictions are unlikely; this is exacerbated in the case of MA by its inherently stochastic nature. In addition, models require input 'data'; material properties (often not known to precision, particularly at the high strain levels to which powder is often subjected during MA) and process characteristics (collision frequency and velocity, also seldom known a priori). Thus, realistic goals of process modeling are to correctly predict general trends, and perhaps even to predict resulting properties/dimensions within an order of magnitude. The benefits of successful models lie not in their abilities to predict outcomes, but in that they
76
C. Suryanarayana
Table 4.4 Room temperature and elevated temperature mechanical properties of commercial ODS nickel- and iron-base superalloys Alloy
MA 6000
MA 754
MA956
Test
0.2% YS
UTS
temperature o f (~
MPa
MPa
% El.
% RA
RT
1220
1253
7.2
6.5
871 (1600)
675
701
2.2
4.6
982 (1800)
445
460
2.8
1.9
RT
586
965
21
33
871 (1600)
214
248
31
58
982 (1800)
169
190
18
34
RT
517
655
20
35
97 69
100 72
12
30
1000 (1832) 1100 (2192)
MA 6000:15 Cr, 4.5 A1, 0.5 Ti, 2.0 Mo, 4.0 W, 2.0 Ta, 1.1 Y203, balance Ni. MA 754:20 Cr, 0.3 A1, 0.5 Ti, 0.6 Y203, balance Ni. MA 956:20 Cr, 4.5 AI, 0.5 Ti, 0.5 Y203, balance Fe.
help identify critical process and material variables, and reduce the amount of testing needed for process optimization".
4.9. I N D U S T R I A L A P P L I C A T I O N S
The major industrial applications of mechanically alloyed materials have been in the areas of thermal processing, glass processing, energy production, aerospace, and other industries. These applications are based on the oxide-dispersion strengthening effect achieved in mechanically alloyed nickel-, iron- and aluminum-based alloys. Mechanically alloyed materials are strong both at room and elevated temperatures (Table 4.4). The elevated temperature strength of these materials is derived from more than one mechanism. First, the uniform dispersion (with a spacing of the order of 100 nm) of very fine (5-50 nm) oxide particles (commonly used are Y203 (yttria), ThO2 (thoria), and La203 (lanthana)), which are stable at high temperatures, inhibit dislocation motion in the metal matrix and increase the resistance of the alloy to creep deformation. Another function of the dispersoid particles is to inhibit the recovery and recrystallization processes, because of which a very stable large grain size is obtained; these large grains resist grain rotation during high temperature deformation. A stable large grain size can also be obtained by secondary recrystallization mechanisms. Second, the very homogeneous distribution of alloying elements during MA gives both the solid-solution strengthened and precipitation-hardened alloys more stability at elevated temperatures and overall improvement in properties. Mechanically alloyed materials also have excellent oxidation and hot corrosion resistance. The increased resistance to oxidation-sulfidation attack is due to the homogeneous distribution of the alloying elements and the improved scale adherence due to the dispersoid itself (Hack 1987, Fischer et al. 1991).
Mechanical Alloying
77
Temperature, OF 1600
90 ~E
&
-~
50
--~
40
"-~
30
~
20
:3
0 0
o
i
" DS Mar-M200 + Hf
A SC PWA 454 o TD-Nickel
70 60
2200
9MA 6000
8O '
x:"
1900
=,,.,..
10 0 700
800
900
1000
1100
1200
Temperature, ~
Figure 4.7. Comparison of 1000-h specific rupture strength of INCONEL MA6000 with dispersion strengthened Mar-M200 + Hf, TD-Ni, and single-crystal PWA 454. Courtesy of Inco Alloys International.
4.9.1 Nickel-base alloys The most significant advantage of oxide-dispersion strengthened superalloys is the improvement in the stress rupture properties. Figure 4.7 compares the specific rupture strength (strength/density) for a 1,000-h life as a function of temperature for several nickel-base superalloys used for turbine blade applications. Mar-M200 is a nickel-base alloy containing by weight percent 9.0 Cr, 5.0 A1, 2.0 Ti, 12.0 W, 10.0 Co, 1.0 Nb and 1.8 Hf, while PWA454 is a nickel-base alloy containing 10.0 Cr, 5.0 A1, 1.5 Ti, 12.0 Ta, 4.0 W and 5.0 Co, and TD Ni is nickel containing 2.0 wt.% ThO2. It is clear from this figure that the MA 6000 alloy can maintain a given stress for a much longer time than a conventional alloy for similar vane applications. This is mainly due to the benefits of the combined strengthening modes in the mechanically alloyed material. The mechanically alloyed superalloys are considered mainly for three groups of applications--gas turbine vanes, turbine blades, and sheets for use in oxidizing/corrosive atmospheres. The largest use of MA 754 is as vanes and bands for aircraft gas turbine en-
78
C. Suryanarayana
~
R
Figure 4.8. Some typical gas turbine engine components made out of INCONEL alloy MA 754. Courtesy of Inco Alloys International.
gines (Fig. 4.8). For applications requiring good resistance to thermal fatigue, such as gas turbine vanes, Inconel MA 754 is given a strong texture. The majority of the grains are aligned so that their (100) axes are parallel to the principal working direction and along the length of the bar. Such texture results in low modulus of elasticity (149 GPa) in the longitudinal direction. The low modulus improves resistance to thermal fatigue by lowering stresses for given thermal strains. The alloy MA 758 is used in the glass industry for high-temperature components requiring both elevated temperature strength and resistance to extremely corrosive molten glass. It is also used for internal combustion engine components, mainly in critical fuel injection parts. One novel industrial application of alloy MA 754 is a high-temperature atmosphere-circulation fan in a "floating" furnace design being commercialized in Japan. Large rounds are used for the hub and plate material is used for blades of the fan, which operates at temperatures over 1,100~ The MA 6000 alloy is a more complex alloy developed as a blade material for advanced gas turbines. It is used for first and second stage turbine vanes and blades machined form solid bar. Unlike cast alloys, MA 6000 exhibits nearly flat rupture-life curves at high temperatures due to the combination of oxide dispersion strengthening and high grainto-width ratios (typically > 10 to 1). Because of its composition, MA 6000 has excellent resistance to oxidation and sulfidation. The characteristics of this alloy allow blade cooling to be reduced or eliminated as the metal temperature can be increased by 100~ or more in engines where the stresses are medium or low.
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79
4.9.2 Iron-base alloys The mechanically alloyed iron-base superalloys combine the high-temperature strength and stability of oxide dispersion strengthening with excellent resistance to oxidation, carburization, and hot corrosion. These alloys are suitable for use in gas turbine combustion chambers. Incoloy alloy MA 956 (20 Cr, 4.5 A1, 0.5 Ti, 0.5 Y203, balance Fe) is particularly well suited for use in heat processing applications. For example, vacuum furnace fixtures made of MA 956 have shown excellent durability and are able to compete with wrought molybdenum, which is also used in these applications. In comparison to molybdenum, MA 956 is about 30% lower in density, providing weight savings and cost advantages. Furthermore, since MA 956 has a lower vapor pressure than molybdenum, it will not coat the inside of the vacuum chamber or the parts being heat treated. Thus, MA 956 rods, flats, and sheets are used in numerous atmosphere and vacuum furnace applications including muffles, baskets, trays, and thermowells. Alloy MA 956 in tubing form has also been used for high temperature, severe service applications such as coke injection lance pipes in steel making. The alloy MA 956 is also being used in glass processing industry because of its resistance to attack by molten glass. Because of this corrosion resistance, the alloy is being evaluated for applications such as firing-kiln rollers, muffle tubes, and furnace racks. Other applications include molten-glass resistance heaters, thermocouple protection tubes, glass-processing components used in nuclear waste disposal and the bushings used to make single and multistrand fibers. More recently, MA 957 (14 Cr, 1.0 Ti, 0.3 Mo, 0.25 Y203, balance Fe) has been evaluated for use as the fuel cladding in fast neutron, breeder reactors. Conventional austenitic alloys are unsuitable for this application due to the dimensional swelling phenomenon caused by the high neutron fluxes. The mechanically alloyed materials are also being evaluated for heat exchanger components in high-temperature gas-cooled reactors.
4.9.3 Aluminum-base alloys The success of mechanically alloyed superalloys led to the development of dispersionstrengthened aluminum alloys. Since an aluminum oxide layer is always present either on the surface of the powder particles at the start of processing or during milling, its incorporation into the alloy contributes to significant improvements in the properties of the alloy. Furthermore, since aluminum is a ductile metal, PCAs are added to assist in minimizing cold welding during processing. Aluminum carbides are formed during MA by the decomposition of the PCA. Both the oxide or carbide type dispersions are about 30-50 nm in size and stabilize the ultrafine grain size. This results in 50% increase in strength, higher fracture toughness and improved resistance to stress corrosion cracking and fatigue crack growth of the mechanically alloyed materials. IncoMAP alloy AL-9052 (4.0 Mg, 1.1 C, 0.6 O, balance A1) has a density 5% less than that of conventional agehardenable aluminum alloy of comparable strengths such as 2024. With its combination of lightweight, high strength, and corrosion resistance, IncoMAP alloy AL-9052 is evaluated for aerospace applications where marine corrosion is also a factor. Addition of lithium to mechanically alloyed aluminum alloys has produced an ultra lightweight alloy IncoMAP alloy AL-905XL (4.0 Mg, 1.3 Li, 1.1 C, 0.6 O, balance
80
C. Suryanarayana
A1). Its density is 8% lower and stiffness 10% greater than the age-hardenable conventional alloy 7075-T73 of comparable strength. The excellent combination of the properties makes this alloy very attractive for airframe applications. In particular, the freedom from age-hardening treatments makes it possible to produce forgings and heavy sections with homogeneous metallurgical structures.
4.9.4 Magnesium-base alloys A useful application of the MA technique was in the production of supercorroding magnesium alloys that operate as short-circuited galvanic cells to corrode (react) rapidly and predictably with an electrolyte such as seawater to produce heat and hydrogen gas (Black 1979). Such an alloy system is suitable as a heat source for warming deep-sea divers, as a gas generator to provide gas for buoyancy, or as a fuel in hydrogen engines or fuel cells. Magnesium-base alloys containing Fe, Cu, Co, Cr, or Ti have been evaluated for such applications. The Mg-5 to 20 at.% Fe alloy is ideal because of its extremely fast reaction rate, high power output, and the high percentage of theoretical completion of the actual reaction. For corrodable release links an alloy with a slower reaction rate, such as Mg-5 at.% Ti is useful. There have also been a number of investigations in recent years to use the MA processing technique to produce metal hydrides. This is because metal hydrides are materials for safe storage of hydrogen and they can store hydrogen with a higher volume density than liquid hydrogen. However, these are sensitive to surface oxidation and hence can be a limiting factor in their commercial utilization. Nanocrystalline hydrides have a high density of defects and interfaces that could enhance diffusion, and therefore nanocrystalline intermetallics would not require activation treatments at high temperatures and pressures after exposure to air (Tessier et al. 1996). Several magnesium-base and iron-base intermetallics are being evaluated for this application. Figure 4.9 summarizes the applications of mechanically alloyed materials. The range of these applications is continuing to increase.
4.10. C O N C L U D I N G R E M A R K S
Mechanical alloying is a simple, elegant, and useful processing technique that continues to attract the serious attention of researchers. Even though the technique was originally developed to produce ODS superalloys, the synthesis of novel crystalline and amorphous phases has spurred lots of research investigations in recent years. It is estimated that so far about 4,000 research/review papers have been published in this area, with nearly 500 per year during the last 2-3 years. Both equilibrium and non-equilibrium phases and commercially useful materials could be synthesized starting from elemental powders. This also appears to be an economical process with vast potential. One of the greatest advantages of MA is in the synthesis of novel alloys that cannot be made by any other technique, such as alloying of normally immiscible elements. This is because MA is a completely solid-state processing technique and therefore limitations imposed by the phase diagrams do not apply here.
Mechanical Alloying
81
I~ STRENGTHENEDALLOYS ) ~_Ni-, Fe-, AND AI-BASE . . ~
Figure 4.9. Typical applications of the mechanical alloying process.
On an industrial scale, this is an accepted process. The ODS alloys have a higher temperature operating capability at the same stress or increased load-bearing capability at the same temperature than alloys without an oxide dispersion. The ODS alloys continue to find applications in a wide variety of industries. The two production facilities of Inco Alloys International have a combined annual powder capacity approaching 300,000 kg. The yield of the final product varies greatly with the product form and size, but the final wrought capacity is over 200,000 kg. Mechanical alloying involves powders with very small sizes and these should be handled with caution and care. Because of the large surface area, they are highly reactive and can be pyrophoric and can cause health problems when inhaled. Precautions should be taken not to open the powder to atmosphere immediately after milling since this can lead to oxidation of the powders and in some situations they can even catch fire.
REFERENCES
Abdellaoui, M. and Gaffet, E. (1996) Acta Mater., 44, 725. Anantharaman, T. R. and Suryanarayana, C. (1987) Rapidly Solidified Metals: A Technological Overview (TransTech Publ., Aedermannsdorf, Switzerland). Arzt, E. and Schultz, L. (1989) New Materials by Mechanical Alloying Techniques (Deutsche Gesellschaft ftir Metallkunde, Oberursel, Germany). Asahi, N., Maki, T., Matsumoto, S. and Sawai, T. (1994) Mater. Sci. Eng., A181/182, 841. Atzmon, M. (1990) Phys. Rev. Lett., 64, 487.
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Avvakumov, E. G. (1986) Mechanical Methods of Activation of Chemical Processes (Nauka, Novosibirsk, Russia). Bahadur, D., Singh, K. and Roy, M. (1992) Mater Sci. Eng., A154, 79. Bahadur, D., Dunlap, R. A. and Foldeaki, M. (1996) J. Alloys & Compounds, 240, 278. Bakker, H. (1998) Mater Sci. Foundations, 1, 1. Bakker, H., Zhou, G. E and Yang, H. (1995a) Prog. Mater Sci., 39, 159. Bakker, H., Zhou, G. E and Yang, H. (1995b) Mater Sci. Forum, 179-181, 47. Beddoes, J. C., Wallace, W. and de Malherbe, M. C. (1992) Inter J. Powder Metall., 28, 313. Beke, D. L., Bakker, H. and Loeff, P. I. (1991 a) Acta Metall. Mater., 39, 1267. Beke, D. L., Loeff, E I. and Bakker, H. (1991b) Acta Metall. Mater., 39, 1259. Benjamin, J. S. (1976) Sci. Amer., 234(5), 40. Benjamin, J. S. (1990) Metal Powder Rept., 45, 122. Birringer, R., Hahn, H., H6fler, H. J., Karch, J. and Gleiter, H. (1988) Defect Diff. Forum, 59, 17. Black, S. A. (1979) Report of Civil Eng. Lab. (Navy, Port Hueneme, CA), 40 pp. Boettinger, W. J. and Perepezko, J. H. (1993) in Rapidly Solidified Alloys: Processes, Structures, Properties, Applications (Marcel Dekker, New York), p. 17. Cahn, R. W. (1991) in Glasses and Amorphous Materials, Vol. 9 of Materials Science and Technology--A Comprehensive Treatment, ed. Zarzycki, J. (VCH Verlagsgesellschaft GmbH, Weinheim, Germany), p. 493. Calka, A. (1991) Appl. Phys. Lett., 59, 1568. Calka, A. and Radlinski, A. P. (1989) Mater Sci. Eng., A l l & 131. Calka, A. and Radlinski, A. P. (1990) J. Less-Common Metals, 161, L23. Calka, A., Nikolov, J. J. and Williams, J. S. (1996) Mater Sci. Forum, 225-227, 527. Chakk, Y., Berger, S., Weiss, B.-Z. and Brook-Levinson, E. (1994) Acta Metall. Mater., 42, 3679. Chattopadhyay, K., Wang, X.-M., Aoki, K. and Masumoto, T. (1996) J. Alloys & Compounds, 232, 224. Chitralekha, J., Raviprasad, K., Gopal, E. S. R. and Chattopadhyay, K. (1995) J. Mater Res., 10, 1897. Cho, Y. S. and Koch, C. C. (1993) J. Alloys & Compounds, 194, 287. Clark, C. R., Suryanarayana, C. and Froes, E H. (1995) in Advances in Powder Metallurgy & Particulate Materials - 1995, Compiled by M. Phillips and J. Porter (MPIE Princeton, NJ), Part I, pp. 1-135. Courtney, T. H. (1994) in Reviews in Particulate Materials, Vol. 2, eds. Bose, A., German, R. M. and Lawley, A. (Metal Powder Industries Federation, Princeton, NJ), p. 63. Courtney, T. H. and Maurice, D. R. (1996) Scripta Mater., 34, 5. Courtney, T. H. and Wang, Z. (1992) Scripta Metall. Mater., 27, 777. deBarbadillo, J. J., Froes, E H. and Schwarz, R. B. (1993) Mechanical Alloying for Structural Applications (ASM International, Materials Park, OH). Desch, E B., Schwarz, R. B. and Nash, E (1991) J. Less-Common Metals, 168, 69. Desch, E B., Schwarz, R. B. and Nash, P. (1996) Scripta Mater., 34, 37. Duwez, E (1967) Trans. Amer Soc. Met. Quart., 60, 607. Duwez, E, Willens, R. H. and Klement, W. (1960) J. Appl. Phys., 31, 1136. Eckert, J., Schultz, L., Hellstern, E. and Urban, K. (1988) J. Appl. Phys., 64, 3224. Eckert, J., Schultz, L. and Urban, K. (1989) Appl. Phys. Lett., 55, 117. Eckert, J., Schultz, L. and Urban, K. (1990a) J. Less-Common Metals, 167, 143. Eckert, J., Schultz, L. and Urban, K. (1990b) Z. Metallkde., 81,862. Eckert, J., Holzer, J. C., Krill, III, C. E. and Johnson, W. L. (1992) J. Mater Res., 7, 1751. E1-Eskandarany, M. S., Aoki, K. and Masumoto, T. (1991 ) J. Japan Soc. Powder Powder Metall., 38(1), 59. Fiorani, D. and Magini, M. (1997), Materials Science Forum, 235-238 (Proceedings of the International Symposium on Metastable, Mechanically Alloyed, and Nanocrystalline Materials, Rome, Italy, May 20-24 1996). Fischer, J. J., deBarbadillo, J. J. and Shaw, M.J. (1991) Heat Treating, 23(5), 15. Froes, E H. and deBarbadillo, J. J. (1990) Structural Applications of Mechanical Alloying (ASM International, Materials Park, OH). Gaffet, E., Malhouroux, N. and Abdellaoui, M. (1993) J. Alloys & Compounds, 194, 339. Gente, C., Oehring, M. and Bormann, R. (1993) Phys. Rev., B48, 13244. Gerasimov, K. B., Gusev, A. A., Ivanov, E. Y. and Boldyrev, V. V. (1991) J. Mater Sci., 26, 2495. Gialanella, S., Newcomb, S. B. and Cahn, R. W. (1992) in Ordering and Disordering in Alloys, ed. Yavari, A. R. (Elsevier, London), p. 67. Gilman, E S. and Benjamin, J. S. (1983) Ann. Rev. Mater Sci., 13, 279.
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Gleiter, H. (1989) Prog. Mater Sci., 33, 223. Goodwin, R S. and Ward-Close, C. M. (1995) Mater Sci. Forum, 17%181, 411. Guo, W., Iasonna, A., Magini, M., Martelli, S. and Padella, E (1994) J. Mater Sci., 29, 2436. Hack, G. A. J. (1987) Metals and Materials, 3, 457. He, L. and Ma, E. (1996) J. Mater Res., 11, 72. Heinicke, G. (1984) Tribochemistry (Akademie Verlag, Berlin). Hellstern, E., Fecht, H. J., Garland, C. and Johnson, W. L. (1989) in Multicomponent Ultrafine Micrtostructures, Vol. 132, eds. McCandlish, L. E., Polk, D. E., Siegel, R. W. and Kear, B. H. (Materials Research Society, Pittsburgh, PA), p. 137. Herr, U. (1995) Key Eng. Mater., 103, 113. Huang, J. Y., Wu, Y. K., He, A. Q. and Ye, H. Q. (1994) NanoStructured Mater., 4, 293. Huang, J. Y., Yu, Y. D., Wu, Y. K., Li, D. X. and Ye, H. Q. (1997) J. Mater. Res., 12, 936. Ivanov, E. (1992) Mater Sci. Forum, 88-90, 475. Ivanov, E., Konstanchuk, I. G., Bokhonov, B. D. and Boldyrev, V. V. (1989a) Doklady Phys. Chem., 304(1-3), 82. Jang, J. S. C. and Koch, C. C. (1990) J. Mater Res., 5, 498. Jangg, G. (1989) in New Materials by Mechanical Alloying Techniques, eds. Arzt, E. and Schultz, L. (Deutsche Gesellschaft ftir Metallkunde, Oberursel, Germany), p. 39. Jangg, G., Kuttner, R and Korb, G. (1975) Aluminium, 51, 641. Klassen, T., Herr, U. and Averback, R. S. (1997) Acta Mater., 45, 2921. Koch, C. C. (1988) Inter Mater Rev., 33, 201. Koch, C. C. (1989) Ann. Rev. Mater Sci., 19, 121. Koch, C. C. (1991) in Processing of Metals and Alloys, Vol. 15 of Materials Science and Technology--A Comprehensive Treatment, ed. Cahn, R. W. (VCH Verlagsgesellschaft mbH, Weinheim, Germany), p. 193. Koch, C. C. (1993) NanoStructured Mater., 2, 109. Koch, C. C. (1994) Inter J. Mechanochem. Mech. Alloy., 1, 56. Koch, C. C. (1996) Scripta Mater., 34, 21. Koch, C. C. (1997) NanoStructured Mater., 9, 13. Koch, C. C., Cavin, O. B., McKamey, C. G. and Scarbrough, J. O. (1983) Appl. Phys. Lett., 43, 1017. Li, W., Suryanarayana, C. and Froes, F. H. (1995a) in Synthesis and Processing of Lightweight Metallic Materials, eds. Froes, E H., Suryanarayana, C. and Ward-Close, C. M. (TMS, Warrendale, PA), p. 203. Li, W., Suryanarayana, C. and Froes, E H. (1995b) in Advances in Powder Metallurgy and Particulate Materials, Part 1, compiled by Phillips, M. A. and Porter, J. (Metal Powder Industries Federation, Princeton, NJ), p. 145. Liebermann, H. H. (1993) Rapidly Solidified Alloys: Processes, Structures, Properties, Applications (Marcel Dekker, New York). Lu, K. (1996) Mater. Sci. Eng. Rept., R16, 161. Lu, L. and Lai, M. O. (1998) Mechanical Alloying (Kluwer, Norwell, MA). Luton, M. J., Jayanth, C. S., Disko, M. M., Matras, S. and Vallone, J. (1989) in Multicomponent Ultrafine Micrtostructures, Vol. 132, eds. McCandlish, L. E., Polk, D. E., Siegel, R. W. and Kear, B. H. (Materials Research Society, Pittsburgh, PA), p. 79. Ma, E., He, J.-H. and Schilling, R J. (1997) Phys. Rev., B55, 5542. Magini, M. and Iasonna, A. (1995) Mater Trans. JIM, 36, 123. Magini, M., Iasonna, A. and Padella, E (1996) Scripta Mater., 34, 13. Martelli, S., Mazzone, G., Scaglione, S. and Vittori, M. (1988) J. Less-Common Metals, 145, 261. Massalski, T. B., Okamoto, H., Subramanian, R R. and Kacprzak, L., eds. (1990) Binary Alloy Phase Diagrams, second edition, (ASM International, Materials Park, OH). Maurice, D. R. and Courtney, T. H. (1990) Metall. Trans., 21A, 289. McCormick, R G. (1995) Mater Trans. JIM, 36, 161. McDermott, B. T. and Koch, C. C. (1986) Scripta Metall., 20, 669. Mizutani, U. and Lee, C. H. (1990) J. Mater Sci., 25, 399. Moore, J. J. and Feng, H. J. (1995) Prog. Mater Sci., 39, 243. Mukhopadhyay, D. K., Suryanarayana, C. and Froes, E H. (1994) Scripta Metall. Mater., 31, 333. Mukhopadhyay, D. K., Prisbrey, K. A., Suryanarayana, C. and Froes, E H. (1996) in Proceedings of the International Conferene on Tungsten and Refractory Metals 3, eds. Bose, A. and Dowding, R. J. (Metal Powder Industries Federation, Princeton, NJ), p. 239.
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C. S u r y a n a r a y a n a
Murty, B. S. and Ranganathan, S. (1998) Inter Mater Rev., 43 (in press). Murty, B. S., Mohan Rao, M. and Ranganathan, S. (1993) NanoStructured Mater., 3, 459. Nasu, S., Miglierini, M., Ishihara, K. N. and Shingu, P. H. (1992) J. Phys. Soc. Japan, 61, 3766. Oleszak, D., Burzynska-Szyszko, M. and Matyja, H. (1993) J. Mater Sci. Lett., 12, 3. Pathak, D., Yamada, K. and Koch, C. C. (1993) in Mechanical Alloying for Structural Applications eds. deBarbadillo, J. J. et al. (ASM International, Materials Park, OH), p. 205. Politis, C., Krauss, W., Leitz, H. and Schommers, W. (1989) Modern Phys. Lett., 3, 615. Radlinski, A. P. and Calka, A. (1991) Mater Sci. Eng., A134, 1376. Rawers, J., Govier, D. and Cook, D. (1995) Scripta Metall. Mater., 32, 1319. Russell, K. C. (1985) Prog. Mater Sci., 28, 229. Saji, S., Abe, S. and Matsumoto, K. (1992) Mater. Sci. Forum, 88-90, 367. Schaffer, G. B. and McCormick, P. G. (1989) Appl. Phys. Lett., 55, 45. Schaffer, G. B. and McCormick, P. G. (1992) Mater Forum, 16, 91. Schropf, H., Kuhrt, C., Arzt, E. and Schultz, L. (1994) Scripta Metall. Mater., 30, 1569. Schulz, R. (1996) Materials Science Forum, 225-227 (Proceedings of the International Symposium on Metastable, Mechanically Alloyed, and Nanocrystalline Materials, Quebec, Canada, 1995). Schwarz, R. B. and Johnson, W. L. (1983) Phys. Rev. Lett., 51, 415. Schwarz, R. B., Petrich, R. R. and Saw, C. K. (1985) J. Non-Cryst. Solids, 76, 281. Seki, Y. and Johnson, W. L. (1990) in Solid State Powder Processing, eds. Clauer, A. H. and de Barbadillo, J. J. (TMS, Warrendale, PA), p. 287. Shingu, P. H. (1992) Materials Science Forum, 88--90 (Proceedings of the International Symposium on Mechanical Alloying, Kyoto, Japan, May 1991). Siegel, R. W. (1991) in Processing of Metals and Alloys, Vol. 15 of Materials Science and Technology--A Comprehensive Treatment, ed. Cahn, R.W. (VCH Verlagsgesellschaft mbH, Weinheim, Germany), p. 583. Singer, R. E, Oliver, W. C. and Nix, W. D. (1980) Metall. Trans., llA, 1895. Stoloff, N. S. and Davies, R. G. (1966) Prog. Mater Sci., 13, 77. Suryanarayana, C. (1995a) Bibliography on Mechanical Alloying and Milling (Cambridge International Science Publishing, Cambridge, UK). Suryanarayana, C. (1995b)Intermetallics, 3, 153. Suryanarayana, C. (1995c) Inter Mater Rev., 40, 41. Suryanarayana, C. (1996) Metals & Materials, 2, 195. Suryanarayana, C. (1998) in Powder Metal Technologies and Applications, ASM Handbook, Vol. 7 (ASM International, Materials Park, OH) (in press). Suryanarayana, C. (1999) Prog. Mater Sci. (in press). Suryanarayana, C. and Froes, E H. (1990) J. Mater Res., 5, 1880. Suryanarayana, C. and Jones, H. (1988) Inter J. Rapid Solidification, 3, 253. Suryanarayana, C., Zhou, E., Peng, Z. and Froes, E H. (1994a) Scripta Metall. Mater., 30, 781. Suryanarayana, C., Li, W. and Froes, E H. (1994b) Scripta Metall. Mater., 31, 1465. Suryanarayana, C., Korth, G. E. and Froes, E H. (1997) Metall. Mater Trans., 28A, 293. Suryanarayana, C., Ivanov, E., Noufi, R., Contreras, M. A. and Moore, J.J. (1999) J. Mater Res., 14 (in press). Takeuchi, T., Yamada, Y., Fukunaga, T. and Mizutani, U. (1993) J. Japan Soc. Powder Powder Metall., 40, 951. Tessier, P., Zaluski, L., Zaluska, A., Str6m-Olsen, J. O. and Schulz, R. (1996) Mater Sci. Forum, 225-227, 869. Trudeau, M. L. (1994) Appl. Phys. Lett., 64, 3661. Trudeau, M. L., Schulz, R., Dussault, D. and Van Neste, A. (1990) Phys. Rev. Lett., 64, 99. Wang, G. M., Campbell, S. J., Calka, A. and Kaczmarek, W. A. (1997) J. Mater Sci., 32,1461. Weeber, A. W. and Bakker, H. (1988) Physica B, 153, 93. Yamada, K. and Koch, C. C. (1993) J. Mater Res., 8, 1317. Yang, E., Wagner, C. N. J. and Boldrick, M. S. (1993) Key Eng. Mater., 81-83, 663. Yavari, A. R. (1995) Materials Science Forum, 179-181 (Proceedings of the International Symposium on Metastable, Mechanically Alloyed, and Nanocrystalline Materials, Grenoble, France, June 27-July 1, 1994). Yavari, A. R. and Desr6, R (1990) Journal de Physique 51, Colloque C4, Supplement No. 14 (Proceedings of the International Symposium on Multilayer Amorphisation by Solid-State Reaction and Mechanical Alloying, Grenoble, France, February 21-23 1990). Yermakov, A. E., Yurchikov, E. E. and Barinov, V. A. (1981) Phys. Met. Metallogr., 52(6), 50. Yermakov, A. E., Barinov, V. A. and Yurchikov, E. E. (1982) Phys. Met. Metallogr., 54(5), 90.
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Yvon, E J. and Schwarz, R. B. (1993) in Structural Applications of Mechanical Alloying, eds. deBarbadillo, J.J. et al. (ASM International, Materials Park, OH), p. 393. Zhang, H. (1993) J. Phys.: Condens. Matter, 5, L337. Zhang, E X. and Wang, W. K. (1996) J. Alloys & Compounds, 240, 256. Zhou, E., Suryanarayana, C. and Froes, E H. (1995) Mater. Lett., 23, 27. Zhou, G. E and Bakker, H. (1996) Scripta Mater., 34, 29.
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Chapter 5 Laser Processing 5.1. Principles of Lasers 5.2. Classifications of Laser Processing 5.3. Analysis of the Laser Melting and Quenching Process 5.3.1 Heat Transfer 5.3.1.1 Absorption Phenomena of Laser 5.3.1.2 Pulsed Laser 5.3.1.3 Continuous Wave (CW) Laser 5.3.2 Kinetic Conditions for Solidification 5.3.2.1 Kinetic Conditions for Amorphous Phase Formation 5.3.2.2 Solidification Modes in the Laser Processing of Materials 5.3.3 Microstructure Selection Maps (MSM) for Alloys 5.4. Laser-Quenching 5.4.1 Amorphous Phase Formation 5.4.2 Formation of Crystalline Phases 5.5. Laser Surface-Alloying and Cladding 5.5.1 Ferrous Alloys 5.5.2 Non-Ferrous Alloys 5.6. Laser-Annealing 5.7. Laser-Beam Joining 5.7.1 Welding of Structural Materials 5.7.1.1 Ferrous-Based Alloys 5.7.1.2 Non-Ferrous Alloys 5.7.2 Microjoining 5.8. Conclusions References
89 89 92 92 92 93 94 96 96 97 97 98 98 99 100 100 101 102 106 106 106 108 112 115 115
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Chapter 5 Laser Processing KOJIRO E KOBAYASHI
5.1. P R I N C I P L E S O F L A S E R S
The laser (light amplification by stimulated emission of radiation) is a unique source of electromagnetic radiation. Lasers are distinguished from other light sources by their property to emit a coherent radiation of high spectral purity (Bimberg 1986). The total power output of a laser can be focused down to a spot of an extremely small diameter and so large intensities can be obtained. The use of focused laser energy allows one to accomplish the heat treatment of materials in a controlled spatial and temporal domain. There are at present a large number of commercially available laser sources that could be used to induce the melting needed for surface alloying. Of these a relatively small group, including the carbon dioxide (CO2), neodymium-doped yttrium-aluminum garnet (Nd:YAG), and EXCIMER (excited dimer), account for nearly all the practically used sources (Draper and Poate 1985, von Allmen 1983). The laser characteristics such as wavelength, output power, beam diameter, output mode, pulse length, and repetition rate are important processing factors. Table 5.1 lists the powers and wavelengths of different lasers. The laser processing system for a pulsed or Q-switched source offers cooling rates exceeding those of rapid solidification from the melt by several orders of magnitude. Qswitched mode operation can provide extremely higher peak power pulses and higher pulse rates than normal pulsed mode. Each individual focused laser event produces a melt spot and those melt spots must be raster-scanned on the surface if area coverage larger than the spot itself is desired. In the case of the continuous wave laser the area per unit time scanned is dependent on the effective spot size, the degree of overlap, and the relative transverse velocity.
5.2. C L A S S I F I C A T I O N S O F L A S E R P R O C E S S I N G
Figure 5.1 shows schematically the major laser processes that relate to the formation of non-equilibrium phases. The definition and characteristics of each of these processes is outlined below (Roessler 1986). Laser surface melting and quenching (Fig. 5.1(a)): A laser-beam is scanned over the surface of a substrate to be melted. Behind the laser-beam the molten pool cools very rapidly because of high heat conduction from the molten pool into the still cold substrate. 89
K. F. Kobayashi
90
Laser beam Molten / pool ~ / .............
~ ~
Laser beam /\Molten Alloying / ~pool powder / ~/
Rapidly solidified layer
(a) Laser surface melting and quenching
(b) Laser surface alloying Laser beam
Laser beam Cladding / powder [ \ [ ~,~,~~,~; 9
Alloyed layer
/ Molten ~ pool ~,/
Clad layer
Region heated above annealing / temperature ~
(c) Laser surface cladding Focused IA~mrhmam
Annealed and
/ rapidlycooled layer
(d) Laser annealing
Weld bead
(f) Laser beam soldering
(e) Laser beam welding l-- ......substrate .....
Laser beam ~ A ~ ~
p
I
~Clusters Plasma plume
(g) Laser ablation
Figure5.1. Schematicrepresentationsofvariouslaserprocessesformaterials.
Laser Processing
91
Table 5.1. Powersand wavelengths of different lasers. Source c02 laser (gas) YAG laser (solid: Nd) EXCIMER (gas: ArF) (gas: KrF)
Wavelength (microns)
Power
10.6 1.06
~50 kW CW: ~6 kW Pulse: ~20 kW ~50 W ~150 W
0.193 0.248
This produces a rapidly solidified surface layer and thereby results in the formation of amorphous or metastable phases on the surface. This process is known as laser-glazing. Laser surface-alloying (Fig. 5.1(b)): In surface alloying with a laser, an alloying material, which is delivered to the surface of a substrate by either preplacing a powder on the substrate or blowing the powder into the laser-generated melt pool, is fused to and mixed with the substrate so that the surface results in an alloyed layer after solidification. The alloyed layer shows a fine and homogeneous microstructure with minimal segregation because of the rapid solidification. The thickness of the alloyed layer can be from a few microns to about 1 mm. Laser surface-cladding (Fig. 5.1(c)): Laser surface cladding is similar to laser surface alloying except that the surface clad layer is fused to the substrate with the minimum of dilution with the substrate. Thus, the objective in this process is to provide the characteristics of a cladding material itself to the surface of a substrate. Clad layers having a thickness of up to several millimeters can be formed. Laser-annealing (Fig. 5.1(d)): A laser-beam is scanned over the surface of a substrate and heats it above an annealing temperature before rapid cooling. The objective of laser annealing is to achieve homogenization, solutionizing, recrystallization and/or removing defects locally in the surface layer. Laser annealing can allow very selective annealing of the substrate and extremely high cooling rates in comparison with conventional furnace annealing. Laser-beam welding (Fig. 5.1(e)): Fusion welding with a laser is a deep-penetration welding process. This process is characterized by the formation of a "keyhole". The laser energy causes local vaporization of material and generates a keyhole filled with vaporized metal through the thickness of the workpiece (Fig. 5.1(e)). As the focused laser-beam is scanned over the weld line, the keyhole surrounded by a molten pool moves. Continuous melting and vaporization occur at the leading edge of the keyhole. The molten metal flows back around the keyhole and resolidifies at the rear edge of the molten pool. Thus, a fusion weld joint with deep-penetration can be obtained. The major advantage of laser welding is selective heat input with narrow fusion and heat-affected zones. On the other hand, rapid heating and cooling cycles may result in the formation of unusual metastable phases in the weld fusion and heat affected zones. So, the effect of microstructural changes on the properties of the weld joint should be considered. Laser-beam soldering (Fig. 5.1(f)): Laser-beam joining can be applied to microjoining in electronics industries. Figure 5.1 (f) shows schematically laser-beam soldering between
K. E Kobayashi
92
the lead of a quad-flat pack surface-mounted component and a pad of a printed wiring board. A laser is irradiated onto the lead and the absorbed laser energy melts a solder. The solder solidifies and the solder joint is made after the laser irradiation is stopped. Laser power, beam diameter, and duration time are precisely controlled. Laser-ablation (Fig. 5.1(g)): Laser ablation is a process for the deposition of thin films. When a high energy laser-beam is irradiated onto a target surface, atoms and ions are emitted from the target to form a plasma plume. The plasma expands, cools and is condensed into clusters. The atoms, ions, clusters and/or molecules are collected on a substrate, resulting in a thin film deposition on the substrate surface. The deposited film usually has a non-equilibrium amorphous structure. The applications and characteristics of the above laser processes will be described in detail below in Sections 5.4 to 5.7.
5.3. A N A L Y S I S OF T H E L A S E R M E L T I N G AND Q U E N C H I N G P R O C E S S
5.3.1 Heat transfer The recent availability of lasers with high-power directed energy sources has helped in the development of rapid melting and solidification techniques in which a semi-infinite bulk substrate in intimate contact with a molten layer acts as the quenching medium. All mathematical treatments of classical and laser heat treatments are based on the fundamental equation for heat conduction:
OT = OT { 8t
)~ OT
)~ OT ~
O ~ C p p Ox ! Cpp 8y
)~ OT~
Cpp Oz)
qi
Cpp
(5.1)
Here T stands for the absolute temperature, t for the time, x, y and z are the coordinates, ~. the thermal conductivity, C p the specific heat, p the density and q i is the rate at which heat is supplied per unit time and volume. If all the parameters involved are known, Eq. (5.1) can be solved numerically by either finite difference or finite element methods (Mehrabian et al. 1979, Bergmann et al. 1981, Fritsch 1986).
5.3.1.1 Absorption phenomena of laser Absorption of laser power by the substrate is related to the laser-solid interaction. When an electromagnetic wave interacts with the atoms in the surface of a substrate, the field causes an electrical current. The Joule effect increases the surface temperature of the substrate due to the finite electrical conductivity of the material. Generally, the power of the laser beam is subdivided into the refected part, the absorbed part and the transmitted part. For metals that are not transparent, the last part can be neglected. Assuming that the Wiedemann-Franz law, i.e., the ratio between the electrical and thermal conductivity is constant, is valid for metals, metals with good conductivity exhibit poor absorption of laser. Similarly, the absorption ratio (fraction of the incident laser energy absorbed by the metal surface) increases with increasing temperature because the conductivity of metals decreases. When the surface is heated continuously by a laser, the surface layers start melting which makes the electrical conductivity lower and makes the absorption ratio higher.
Laser Processing
o 0.8
93
Ag
,,m
(u L.
s 0
0.6
~-
0.4
Cu
i n
k,.
0 .Q <
Fe
0.2
0.4
1
2
4
8
W a v e length (l~nn) Figure 5.2. Changein absorption ratio as a function of wavelengthfor metal surfaces.
In electromagnetic fields, the frequency affects the electrical conductivity and hence the laser light with different wavelengths shows different absorption values for a given material. Figure 5.2 shows the change in absorption ratio as a function of laser wavelength for metal surfaces. It is recognized that the CO2 laser with a 10.6 # m wavelength has an inferior absorption ratio to Nd: YAG laser with shorter wavelength (1.06/zm). Additionallyl reactions with the molten metal and the surrounding atmosphere may also affect the absorption ratio.
5.3.1.2 Pulsed laser The nature of interaction between a laser-beam and a material surface is controlled by two variables. One is the absorbed power density from the laser and the other is the interaction time. In the pulsed mode, interaction time means the time of laser irradiation. In the rapid solidification surface treatment process, glazing requires a high power density and substantially shorter interaction times. These requirements are met by the application of an Nd: YAG laser with Q switch. As discussed in detail by Bloembergen (1979) and others (Spaepen and Turnbull 1982, Linet al. (1984), picosecond pulse lasers can be used to induce very high quench rates, up to 1012 K/s, in thin molten overlays on a metallic substrate. During laser-glazing, the modified surface is considered to go through the following stages: (i) energy deposition in a layer on the order of the absorption depth (~50 nm), (ii) very fast melting ( ~ 1000 m/s) of a ~ 100 nm thick layer, and (iii) cooling and solidification.
94
K. F. Kobayashi
The quenching rate, u, is determined as a function of temperature gradient at the interface from the thermal balance.
u --
(
~.V 6 Hf
0T) Ox interface
9
(5.2)
o
Here V and A H f represent the molar volume and molar heat of fusion, respectively. On the other hand from the kinetics of interface motion u = u
Tm-T/ ~ Ti
0
(5.3)
where Tm stands for melting temperature, T/ for interface temperature and u0 is the velocity prefactor. For collision-limited motion u0 almost equals to the speed of sound, while for diffusion-limited motion uo = D~ I, where D is the diffusivity in the liquid and 1 is the interatomic distance. Generally u0 has a value between these two extremes because of a mixed mechanism. Since the temperature gradients in this process are very steep (1011 K/m during melting and 101~ K/m during cooling), u from Eq. (5.2) becomes very large and at the same time, Ti becomes considerably different from Tm from Eq. (5.3). The lifetime of the melt, determined experimentally (Lin and Spaepen 1981, Lin and Spaepen 1983) or estimated from heat flow calculations, is reported to be about 1 ns. Since the temperature of the material rises several thousand degrees in that period, the cooling rate can be estimated to be at least d T / d t = 103 K/10 -9 s = 1012 K/s (Lin and Spaepen 1984). This rapid solidification can lead to the formation of new metastable crystalline phases, or, if the crystal growth kinetics are slow enough, glass formation results. 5.3.1.3 Continuous wave (CW) laser In the pulsed mode, very high cooling rates are achieved due to the short interaction time. At the same time, however, this makes the modified region extremely small. Normally for industrial purposes, very high cooling rates are not required and the continuous mode is applied to modify the surface effectively 9 The absorbed heat-flux distribution, which is a function of the position and time, is treated moving in a CW laser while it is treated as stationary in a pulsed laser. In the basic investigation of laser-beam processing, an important aspect is the use of appropriate heat flow models. If all the parameters in Eq. (5.1) are known, it can be solved numerically by either finite difference or finite element methods 9 However, the material parameters are not always constant and their temperature dependence is not always known. Internal heat sources or sinks during transformation, melting or evaporation are not also always negligible. Cline and Anthony (1977) showed that the moving point source model, which involves an analytical solution of the thermal diffusion equation using Green's functions, could be used to estimate the temperature distribution during CW laser irradiation. During exposure and removal of a laser with a uniform energy flux onto the surface, the heating and cooling of a semi-infinite plate during the glazing process can be modeled conveniently and simply (Greenwald et al. 1979).
95
Laser Processing
During heating, the temperature at a point x below the surface at time t is given by (5.4)
T ( x , t) -- To + qJ(x, t)
and 2qi[~t 7*(x,t)----~
(
-exp
where To is the ambient temperature and tc During cooling
-
-
-~
x2 )
-
x err ( 2 ~ x t ) ]
(5.5)
)~/(CplO) is the thermal diffusivity.
T ( x , t) -- To + qJ(x, t) - ~P(x, t - r)
(5.6)
where r is the interaction time, defined in continuous mode as (beam diameter)/(beam traveling speed). The times for initial melting, r* and final solidification r** are given by Eqs. (5.7) and (5.8). r* =
Tm-T0
ab
(5.7)
(r + r*) 2 r** =
4r*
(5.8)
Here a and b are defined as follows: a -- v/rr/4K
(5.9)
b - qi/)~.
(5.1 O)
From these relations, the trailing liquid zone length, L, may be determined, since L is simply equal to v(r** - r*), in which v is the beam traveling speed. The compatibility of the model is checked by comparing this calculated length with the experimental value. Under conditions in which the energy inputs range from 1 to 100 J/cm, the heating effect is concentrated in a very thin region at the material surface. This gives rise to extremely high thermal gradients, which promote rapid solidification of the melt. At the same time, the melt and the substrate are in intimate contact, which makes the heattransfer coefficient between them infinity. These effects enable the solidification to occur faster in the case of atomization and splat-cooling types of processes. As a consequence of the extremely rapid solidification, a variety of interesting metallurgical microstructures can be developed, some of which are unique to the laser glazing process.
96
K. E Kobayashi
5.3.2 Kinetic conditions f o r solidification Crystallization kinetics, cooling rate, and laser parameters are closely connected for the intelligent processing of materials. 5.3.2.1 Kinetic conditions for amorphous phase formation The conditions for amorphous phase formation are to be derived from the nucleation and growth theory (Uhlman 1972, Davies et al. 1974, Kobayashi and Shingu 1988). Classical nucleation theory gives the following expression for the nucleation frequency J (T) (Turnbull 1950, Pound and Lamer 1974, Miyazawa and Pound 1974, Spaepen 1975, Spaepen and Turnbull 1976):
100 [
J ( T ) -- ~ e x p rl
-
3
1 - Tr
(5.11)
where r/= viscosity of liquid, Tr = reduced temperature (=T/Tm) and Tm = equilibrium melting point; c~ is a factor which depends on the atomic arrangement at the interface with a value close to unity, and/3 is the entropy change due to solidification in units of the gas constant. The growth rate of crystal nuclei is given by molecular kinetic considerations assuming that the atomic transport across the interface region, which usually is assumed as several atoms diameters in thickness, takes place by a diffusive process according to the equation: Ri(T) :
K r/(1 -- Tr)
(5.12)
where K = d 3L f/32rq6 and d = atomic diameter, L f = heat of fusion, and q~ = interface thickness. The temperature dependence of viscosity in the undercooled liquid region is represented assuming a Tamrnann and Hesse type of relation, as log 7? = A +
B
T - T~
(5.13)
where A, B and T~ are constants which are determined assuming r/ -- 1013 poise (10 le Pa sec) at T -- Tg (glass temperature); r/ is the value for the Arrhenius relation at T = Tm and d(log r/)/d(1/T) is the slope of the Arrhenius relation at T = Tm. The general form of the fraction of phase transformed X may be given as - In(1 -- X) -- -4n. ~ fot J (T) [fo t Ri (co) do)]3 dO
(5.14)
where J ( T ) is the nucleation frequency, Ri(w) is the linear growth velocity, X the crystallized fraction and t the time after passing the melting point during cooling. The temperature-time-transformation (TTT) diagram and the continuous-coolingtransformation (CCT) diagram can be obtained from Eq. (5.14) (Cahn 1956, Grange and Kiefer 1941). By the use of Eq. (5.14) we can estimate the condition for amorphous
Laser Processing
97
phase formation by requiring that the crystallized volume fraction X in the case of laser quenching be very low. 5.3.2.2 Solidification modes in the laser processing o f materials
The surface treatment of materials by laser is of great interest in many technical applications. By moving at high speed, a high power beam, such as that produced by a CO2 laser, over a surface heat treatment or rapid melting-solidification of the base material can be obtained at depths varying from a few microns to several millimeters. For the case of CW laser sources and metals, heat flow has been modeled by Mehrabian and co-workers (Mehrabian 1982, Hsu et al. 1978, Hsu et al. 1980, Kou et al. 1981, Sekhar and Mehrabian 1981) and others (Mazumder and Steen 1980, Chande and Mazumder 1982, Chande and Mazumder 1983, Lin and Spaepen 1984). During laser remelting, the rate of solidification governs the possible formation of metastable phases, their composition, and the scale of the microstructure (Boettinger et al. 1980, Kurz et al. 1986, Trivedi et al. 1987). Laser resolidification starts epitaxially on a crystalline substrate and nucleation plays a less important role. At the high rates of absolute stability, one cannot assume equilibrium partitioning any more. Such departure from equilibrium partitioning was discussed by Boettinger et al. (1984) who used two response functions to describe the temperature and the solid composition at the interface, where the variation of the partition coefficient with growth velocity was given by Aziz (1982) and Jackson et al. (1980). An amorphous phase should form in two different cases. In the first case, the cooling rate is too fast for partitioning of atoms to occur. In the second, partitioning and/or rearrangement of atoms for crystallization is too slow for a given cooling rate (Kobayashi and Shingu 1982, Kobayashi et al. 1984). 5.3.3 Microstructure selection maps ( M S M ) f o r alloys
Recently a series of solidification models have been developed which are valid at low and high growth velocities for different microstructures (Trivedi et al. 1987, Kurz et al. 1986, Boettinger et al. 1988, Lipton et al. 1987, Carrad et al. 1992). These models predict phase selection and the solidification microstructure of an alloy as a function of solidification parameters such as the growth rate and temperature gradient. These are discussed for dendrites (Kurz et al. 1986, Gill and Kurz 1995, Gilgien et al. 1995, Kurz 1995) or eutectic (including peritectic) structures (Rappaz et al. 1987, Kurz and Trivedi 1994, Gill and Kurz 1995, Gilgien et al. 1995) in rapid solidification processing such as in welding and laser treatment. MSMs contain information on the phase selection, solidification morphology and sometimes size of the microstructure as a function of composition and growth rate (or growth temperature, often combined with the equilibrium phase diagram). The modification of laser treatment conditions by controlling the atmosphere and the initial temperature of the sample allowed the reduction of instabilities in the liquid bath at low beam velocities. This made it possible to observe the metastable eutectic A1-A16Fe and the transition from the metastable eutectic to Al-dendrites for A1-4at%Fe alloy during laser resolidification (Gilgien et al. 1995). In the case of stainless steel, metastable phases such as y may form in a composition range where under normal welding conditions
98
K. F. Kobayashi
still forms. This becomes clear when one considers the fact that a phase with less solute rejection is preferred at high growth velocities even if its liquidus temperature is lower than that of the equilibrium phase (David and Vitek 1989). In high power laser treatment (>5 kW) or in welding, the melt pool might become so large that buoyancy driven convection becomes important (Kurz and Trivedi 1994). Preheating lowers the surface temperature gradient that decreases the driving force for Marangoni convection. Preheating is therefore useful for low velocity laser resolidification.
5.4. LASER-QUENCHING
5.4.1 Amorphousphaseformation Laser quenching yields an amorphous phase over most of the phase diagram in Pd-Si and Pt-Si systems (von Allmen et al. 1980). For the formation of an amorphous phase by laser-alloying methods, it must not only be produced as a result of an individual laser event, but the amorphous region thus produced must remain amorphous throughout subsequent heat cycling needed for overlap in multi track treatment. Fe-B alloys between 5 and 24 at%B were found to be amorphous following laser-alloying (Lin and Spaepen 1982). In this work, Nd : YAG laser melted the alloy for 30 picoseconds (ps) and it was quenched at the high cooling rate of about 101~ K/s. The formation of amorphous phases from pure metallic melts by short laser pulses has also been reported. Frequency-quadrupled Nd-laser pulses (wavelength: 0.265 #m) of 10 ns duration incident on single-crystal Si formed amorphous zones a few tens of nm thick (Tsu et al. 1979). Liu et al. (1979) formed amorphous Si by using 30 ps pulses of both 0.53 and 0.265 #m radiation and the thickness of the zone was estimated to be about 30 nm. Amorphization kinetics of Si prepared by short laser pulses were discussed by Cullis et al. (1982) and by Rozgonyi et al. (1982). In the case of pure A1, the reported amorphous region contained small grains of crystalline A1N (Mazzoldi et al. 1980). Recently, amorphization of Co-Ti thin films was achieved after quenching following melting by a 5 ns Nd: YAG laser pulse (Vitta et al. 1994). The Co-Ti system is suitable for the study of amorphous phase formation at compound-forming compositions since the compounds exhibit a variety of homogeneity ranges and structures. The competition between the formation of an amorphous phase and the terminal solid solution was investigated in this work. The solidification was partitionless and complete amorphous phase formation was found for Ti contents greater than about 23 at% in laser quenched sampies. This compositional limit of the amorphous phase formation agrees well with the prediction based on the influence of atomic size difference on the topographical stability of terminal solid solution that was proposed by Egami and Waseda (1984). Amorphous alloy films were deposited by EXCIMER laser ablation using sintered Ta-Ni targets (Yano et al. 1996). By irradiating a solid target with a pulsed laser-beam of high energy density, atoms and ions forming plumes are emitted from the target, and deposited on a substrate as an amorphous phase. These Ta-Ni amorphous alloy films are known to have extremely high corrosion resistance in both acid and alkaline solutions (Mitsuhashi et al. 1987, Lee et al. 1995). The pulsed laser deposition was performed with
Laser Processing
99
a KrF EXCIMER laser (Lamda Physics LPX305i) having a wavelength and beam dimensions 248 nm and 15 mm x 35 mm, respectively. In another case, the effect of nitrogen on the pulsed laser deposition of amorphous silicon carbide films was investigated to evaluate the properties and structure (Yee et al. 1996). The mechanical, optical, and electrical properties of amorphous carbon-based films (Kurokawa et al. 1987, Fraunheim et al. 1989, Elhossary et al. 1990, Liu and Cohen 1990, Lu et al. 1992, Schwan et al. 1994) can be controlled by focusing laser pulse from ArF EXCIMER laser (Questek 2740SC) operating at 193 nm onto a target. CO2 laser surface treatment consisting of laser surface melting and subsequent selfquenching is also useful to form amorphous phases (Bergmann and Mordike 1981, Kumagai et al. 1986, Hirose et al. 1992, Hirose et al. 1994). An amorphous surface has a variety of useful characteristics such as extremely high corrosion resistance and unique electrocatalytic activity. Hirose et al. (1994) fabricated an amorphous magnetic alloy layer by CO2 laser quenching. The laser-quenched Fe78B 13Si9 alloy layer consisted of an amorphous phase and was utilized to construct a non-contact-type static torque sensor with good sensitivity and linearity.
5. 4.2 Formation o f crystalline p h a s e s
Laser melting and quenching is used not only for the formation of amorphous phases but also for understanding the solidification process. During CO2 laser surface resolidification of Ni3 (A1, Ti) alloys, a sharp transition was observed with increasing growth velocity from the growth of an ordered intermetallic to metastable structures (Hunziker and Kurz 1997). This transition is attributed to disorder trapped in the ordered phase and the corresponding undercooling. High-intensity laser irradiation was used for melting silicon carbide in a high-pressure inert gas atmosphere (Sadler and Shamsuzzoha 1997). By this technique, an entirely new class of materials such as cast silicon carbide shapes and silicon carbide fibers spun or blown directly from the melt might be possible. Badan et al. (1985) formed cellular or dendritic layers, with carbon and chromium supersaturation and metastable austenite as the main constituent on the steel surface by CW-CO2 laser. In the case of YAG laser processing, non-crystalline Nb40Ni60 was formed (Lin and Spaepen 1986) and the morphology of nanostructures on the scale of 1-3 nm could be altered by the laser quenching process (Scandella et al. 1991). Using KrF (248 nm) EXCIMER laser of relatively long (45 ns full width half maximum (FWHM)) pulse duration, melting and subsequent solidification of thick (190410 nm) amorphous-silicon (a-Si) layers produced by ion implantation was investigated (Lowndes et al. 1987). "Fine grained polycrystalline silicon" region produced by explosive crystallization of a-Si contained large numbers of disk-shaped silicon flakes, and the onset of epitaxial regrowth was marked by a slight decrease in surface melt duration. KrF EXCIMER pulsed laser deposition (PLD) has been employed successfully to grow high-quality films that are fully dense, long lasting, and have low friction surfaces. Film properties such as crystallinity, stoichiometry, morphology and dopant concentration can easily be adjusted by varying PLD parameters, since it is exactly these properties that are important for determining the tribological behavior. The feasibility of using PLD to grow
100
K. F. Kobayashi
WS2 films has been evaluated (Prasad et al. 1992). PLD was used to grow adherent films of WS2 on 440C stainless steel substrates (Zabinski et al. 1994). EXCIMER laser is quite a useful tool for the synthesis of electronic materials. A pulsed laser deposition technique formed strained dielectric superlattices of BaTiO3 (BTO) and SrTiO3 (STO) (Tabata et al. 1994). A large stress of 400-500 MPa is introduced at the interface between the BTO and STO layers. A large dielectric constant of 900 was observed with a stacking periodicity of 2 unit cells/2 unit cells. The superlattices show drastically different electrical behavior from that of the solid solution (Sr, Ba)TiO3 films. Dielectric superlattices of (Sr, Ca)TiO3/(Ba, Sr)TiO3 also formed by a pulsed EXCIMER laser ablation (Tabata and Kawai 1997a). The crystal structure is controlled with atomic order accuracy and a large lattice stress of 0.5-1 GPa can be introduced periodically at the interfaces owing to the lattice mismatch of the constituent layers. For the multi-layered films, a laser ablation technique produced a new concept device with a combination of ferroelectric and ferromagnetic materials (Tabata et al. 1997b). An ideal hetero-epitaxy can be obtained owing to the similar crystal structure of perovskite type ferroelectric Pb(Zr, Ti)O3; (so called PZT) and ferromagnetic (La, Sr)MnO3, which is useful for application in DRAM (dynamic random access memory) and SRAM (static random access memory) devices. For the superconducting materials, a laser ablation technique using ArF EXCIMER laser formed Bi-based layered perovskite oxide films (SrBi2Ta207) on the Nb-SrTiO3 substrate or Bi2Sr2CuO6 superconducting bottom electrode with c-axis perpendicular to the substrate surface (Tabata et al. 1995). A multi-target, laser-ablation technique inserted layers of undoped, semiconducting La2CuO4 or SmzCuO4 in a controlled manner in the Lal.ssSr0.15CuO(LSCO) compound superlattices (Tabata et al. 1993). 5.5. L A S E R S U R F A C E - A L L O Y I N G
AND CLADDING
Laser surface treatment is a material-processing method that utilizes the high power density (CO2 laser) available from focused laser sources to melt metal coatings and a portion of the underlying substrate. Laser cladding uses a laser-beam to melt and alloy powders of different composition placed on top of the substrate as well as a thin layer of the substrate. The main purpose of laser cladding is to form a thin interfacial layer of an alloy on a given substrate with minimum dilution of the clad layer and to impart good surface wear and corrosion resistance to the material. The advantages of this technique are production of noble alloys, minimized clad dilution, reduced alloy material loss, reduced machining costs, and reduced distortion. 5.5.1 Ferrous-alloys
A typical use of laser surface treatment is to modify the surface structure of gray cast iron, with the aim of producing a white ledeburitic microstructure while retaining the gray core (Bergmann 1985). The advantage of such a treatment is the combination of good core ductility, a result of stable Fe-C solidification producing an iron-graphite microstructure, and surface wear resistance resulting from the Fe-Fe3C layer obtained in a metastable state.
Laser Processing
101
The phase analysis of the metastable structure was investigated in a laser melted Fe-C alloy (Zarubova et al. 1996). As the laser transformation hardening is characterized by extremely fast heating and self-quenching rates, the presence of alloying elements such as Cr and Mn in steels, that promote hardenability and tend to partition to the carbides in the initial microstructures, is likely to prevent complete austenitization and development of a nearly fully martensitic hardened layer (Bradley and Kim 1988). To improve wear resistance of the AISI 304L austenitic stainless steel, without degrading the inherent corrosion resistance, laser surface treatment was utilized. For this purpose, molybdenum and tantalum were used as alloying elements and they formed a substitutional solid solution in the austenite phase (Akgun and Inal 1995). A 10 kW CO2 laser clad Fe-Cr-Mn-C alloy on an AISI 1016 steel produced a very fine grained microstructure of ferrite and complex carbide precipitation (Singh and Mazumder 1987a). The clad exhibited far superior wear properties compared to Stellite 6 during block-on-cylinder tests. The improved wear resistance was attributed to the fine distribution of metastable M6C carbides in the ferrite matrix. Laser clad Ni-Fe-Cr-A1Hf alloys showed a high degree of grain refinement, considerable increase in solubility of Hf in matrix and the occurrence of Hf-rich precipitates and new metastable phases which contributed towards an increase in the oxidation resistance of the alloy (Singh and Mazumder 1987b). A 10 kW CO2 laser-clad Ni-Cr-A1-Y alloy with either niobium carbide (NbC) or chromium carbides (Cr3C2, Cr7C3, Cr23C6) had higher wear resistance than that of the Ni-Cr-A1-Y alloy cladding without the carbides (Hirose and Kobayashi 1995). Laser cladding of Ni-Cr and Ni-P alloys on steel enhanced the corrosion resistance (Renaud et al. 1991, Garcia-Alonso et al. 1996).
5.5.2 Non-ferrous alloys
Aluminum, titanium, and copper alloys are the most popular materials, but their low wear resistance prevents their widespread applications. A titanium aluminide layer was formed on an aluminum surface by laser processing. By applying a laser-beam to a powder mixture of aluminum and titanium previously placed on the aluminum base metal, the mixture was superficially melted and titanium aluminide alloy layers with a thickness of 400 # m were formed at a relatively high speed. By forming this titanium aluminide alloy layer on an aluminum surface, the wear properties of aluminum have been considerably improved without losing the advantage of lightness of this material (Uenishi et al. 1992a, Uenishi et al. 1992b). A13Ti-based intermetallic compound matrix composite (IMC) layers were formed by using a mixture of aluminum, titanium and ceramic powders as preplaced powders (Uenishi and Kobayashi 1993, Uenishi and Kobayashi 1994, Uenishi and Kobayashi 1996a, Uenishi and Kobayashi 1996b). These IMC layers exhibited good wear resistance. The same technique can be applied to modify the surface of titanium (Hirose et al. 1993). The intermetallic compound clad layers were successfully formed on titanium substrate. A typical example of the clad layers is shown in Fig. 5.3. The clad layers showed good wear and oxidation resistance.
K. E Kobayashi
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Surface modification of copper using a laser is suggested to improve the strength and wear resistance of pure copper at elevated temperatures without sacrificing its electrical properties. In particular, Cu-Cr alloyed layers were produced on a pure copper substrate by irradiating a laser-beam onto chromium powder beds or plasma spray-deposited on the substrate surface (Hirose and Kobayashi 1994).
5.6. L A S E R - A N N E A L I N G
Laser-annealing is characterized by the following advantages in comparison with conventional processes such as furnace annealing" 9 Chemical cleanliness,
Laser Processing
103
9 Extremely short duration of heating cycle and, therefore, high heating and cooling rates, 9 Less thermal damage and distortion, and 9 Controlled shape and location of the treated region. Because of its selective heating of the workpiece, laser-annealing is promising for heat treatment of semiconductor materials as an alternative to furnace-annealing in the microelectronics industry. In conventional IC (integrated circuit) fabrication, the lattice damage caused by ion implantation is removed and the implanted dopants are activated by furnace annealing for at least 30 min at temperatures of 820 K to 1470 K. Laser annealing can more rapidly and effectively remove the lattice damage and activate the dopants. Although significant damage usually remains in the form of dislocation loops after furnace annealing, pulsed-laser annealed specimens have no damage down to a resolution of approximately 1 nm (Young and Wood 1982, Young et al. 1982). Figure 5.4 shows the carrier concentration as a function of implanted dose for pulsed laser and thermally annealed boron-implanted silicon (Young and Wood 1982). In laser-annealed silicon the carrier concentration increases linearly beyond the dose of 1.5 • 1016/cm2 that corresponds to the equilibrium solubility limit of boron in silicon in this case, whereas it saturates for thermal annealing. It is clearly shown in this figure that the non-equilibrium process of laser annealing can allow active carrier concentration in excess of the equilibrium solubility limit. This is achieved due to the extremely high heating and cooling rates in the laser processing. In laser-annealing, the very short duration time and the ease of controlling the shape and the location of the treating region are also beneficial to form shallow junctions, which are needed for highly integrated devices such as megabit SRAMs (Tsukamoto et al. 1992). In conventional furnace annealing or rapid thermal annealing systems, shallow p + n junction formation using boron ion implantation is difficult, because the implanted boron diffuses very rapidly into silicon on subsequent annealing (Tsukamoto et al. 1992). An ultrashallow p + n junction, less than 50 nm deep, has been formed by low power BF2 ion implantation followed by EXCIMER laser-annealing with single pulse irradiation (Tsukamoto et al. 1992, Tsukamoto et al. 1996). Laser-annealing has also been applied to recrystallize amorphous layers of materials to improve their electrical properties. Polycrystalline silicon thin-film transistors (polySi TFTs) are used for large-area electronics applications such as liquid crystal displays and image sensors, because poly-Si TFTs have higher carrier mobility than amorphous silicon (a-Si) TFTs. EXCIMER laser annealing has been applied to the fabrication of highperformance poly-Si TFTs on non-heat-tolerant glass substrates, because its extremely short duration and shallow melt-regrowth avoids thermal damage to the glass substrate (Sameshima et al. 1986, Sera et al. 1989, Kuriyama et al. 1991). Recently large-grain poly-Si TFTs have been obtained by controlling the solidification velocity during dualEXCIMER laser irradiation both from the front side and from the back side (Shimizu et al. 1993) and by low temperature (<673 K) substrate heating during EXCIMER laser annealing (Kuriyama et al. 1992, Kuriyama et al. 1993). Besides the microelectronics applications, annealing of amorphous materials by laserbeams has been applied to improve the magnetic properties of materials (Matteazzi et al.
K. F. Kobayashi
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1985, Sato et al. 1985, Lanotte and Lannotti 1995). Due to the very short duration time and rapid temperature change, laser-heating can produce uniform crystallization at a temperature that is much higher than that reached by other methods (Lanotte and Lannotti 1995). This is very useful for the synthesis of nanocrystalline materials through crystallization of amorphous materials. It has been known that nanocrystalline ferromagnetic materials have improved soft magnetic properties in comparison with crystalline or amorphous materials (Herzer 1992). Pulsed YAG laser irradiation of thick Fe80.5(BSiC)19.5 amorphous ribbons (>50 mm), which is a transformer core material, has decreased the core losses because of refinement of the domain widths (Matteazzi et al. 1985). CW CO2 laser annealing of Fe73.sCulNb3Si13.sB9 amorphous ribbons produced nanocrystalline structure of 50 nm grain size (Lanotte and Lannotti 1995). The laser annealed materials had good soft magnetic properties (Lanotte and Lannotti 1995). It has also been reported that continuous laser-annealing improved the soft magnetic properties of Fe40Ni40P14B6 amorphous ribbons (Sato et al. 1985). Selective and local annealing by laser-beam can be applied to structural materials. Laser irradiation for solution-annealing (laser solution-annealing) has been performed on the surface of a Ni-base superalloy, Inconel 718, to reduce its sensitivity to hydrogen embrittlement (Hirose et al. 1996). Inconel 718, which is one of the most used Ni-base superalloys, is age-hardened by the homogeneous precipitation of yt and y~f phases, which allows the alloy to have a high strength at moderate temperatures, and good creep and fa-
105
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tigue resistance. However, the high sensitivity of Inconel 718 to hydrogen embrittlement is a serious problem in using the alloy, for example, in a liquid hydrogen fueled rocket engine. Hydrogen-induced cracking tends to occur at the surface of stress concentrated regions. Reducing the hardness of such surface regions locally is known to be effective in preventing such cracking. To this end, laser surface solutionizing has been applied to Inconel 718 using a 2.5 kW CO2 laser. The surface-solutionized zones can be obtained without melting the treated zones by optimizing the laser conditions. While the base metal has a hardness of approximately 450 Hv, the hardness of the laser solutionized zones is reduced to below 250 Hv as shown in Fig. 5.5. No precipitates of yl and y1~ phases are identified in the laser solutionized zones by TEM observations. The surface regions are heated beyond the solvus temperatures of y~ and y~ phases and these phases are dissolved in the matrix. The subsequent rapid cooling can prevent the precipitation of the age hardenable phases. Surface-solutionizing can be achieved by laser-scanning. The desired depth of the solutionized zones can be obtained by controlling the laser parameters, traverse speed and defocus distance, as shown in Fig. 5.6. The ductility of the surface solutionized specimen is almost twice that of the base metal in a tensile test under a 29.4 MPa hydrogen atmosphere at room temperature with hydrogen pre-charging. Thus, the laser surface softening is effective in reducing the sensitivity of Inconel 718 to hydrogen embrittlement with little sacrifice to the strength.
106
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Figure 5.6. Effectof traverse speed on depth of the laser surface solutionized zone for different defocus distances (Hirose et al. 1996).
5.7. LASER-BEAM JOINING Joining is one of the most important applications of lasers including both solid-state and gas lasers, ranging from joining and bonding devices in microelectronics to welding structural materials. Laser-beam joining is characterized by its high energy density and flexibility, and has obvious advantages, such as high speed, narrow fusion and heat affected zones (minimizing distortion and damage of adjacent areas), ability of remote noncontact processing, and ease of automation. 5. 7.1 Welding o f structural materials 5. 7.1.1 Ferrous-based alloys Laser-beam welding is an attractive process for joining steels due to its high productivity and possibility of automation. In particular, automotive and shipbuilding industries have an interest in the applications of laser-beam welding. Laser-beam weld fusion and heat affected zones are subjected to significantly more rapid heating and cooling cycles compared to conventional arc welds. In the case of laser-beam welding of C-Mn steels, HY (high yield strength) steels and HSLA (high strength low alloy) steels, metastable phases such as martensite and bainite having high hardness are produced in the weld fusion and heat affected zones (Metzbower and Denney 1986, Moon et al. 1987, Metzbower 1990). Thus, these laser-beam welded zones have higher hardness in comparison with arc welded zones. Therefore, predictions of hardness of laser-beam welds are desired. Ion et al. have developed simplified analytical models of heat flow and phase transformations, and applied them to laser-beam welding of C-Mn steels to predict the maximum heat affected zone hardness (Ion et al. 1996). In this model the heat cycles in the heat affected
107
Laser Processing
Cooling time from 1073K to 773K o~" 0.30 W W r
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Figure 5.7. Diagram of carbon equivalent against absorbed energy showing theoretical contours for maximum heat affected zone hardness in carbon manganese steels (Ion et al. 1996).
zones have been simulated with a formula based on the thin plate analysis of Rosenthal (1946). Microstructural constituents in the heat affected zones have been estimated from the simple carbon equivalent, Ceq = C + Mn/12 + Si/24 (where the concentrations of the different elements are given in wt.%), and cooling time between 1073 K and 773 K using empirical formulas obtained by analyzing experimental data. The hardness of individual phases, namely martensite, bainite and ferrite/pearlite, has been empirically obtained from the carbon equivalent. Finally the maximum heat affected zone hardness has been determined by summing the contributions from the individual phases using a simple rule of mixtures. Figure 5.7 shows the diagram for predicting the maximum heat affected zone hardness from the carbon equivalent and absorbed energy using this model. In laser-beam welding of stainless steels the microstructural changes in weld fusion and heat affected zones due to rapid cooling, which affect both the mechanical and corrosion properties, should be considered. In conventional arc welding, the room-temperature microstructures of weld-bead are estimated using Schaeffler diagram (Schaeffler 1949). However, the solidification microstructures of laser welds subjected to more rapid cooling and higher solidification rate than arc welds are significantly different from those estimated by Schaeffler diagram (Vitek et al. 1982, David and Vitek 1985, David et al. 1987, Khan et al. 1988, Zambon and Bonollo 1994). It has been shown that a significant narrowing of the duplex ferrite-austenite region in the diagram occurs with cooling rates typical of laser-beam welding as schematically shown in Fig. 5.8 (David et al. 1987). Thus, while
K. F. Kobayashi
108 A
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Ferrite (F)
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32 36 40 (mass%)
Figure 5.8. Schaeffler diagram showing schematically the convergence of the 0 and 100% ferrite lines in laser-beam welding (Davidet al. 1987).
austenitic stainless steels such as type 304, 308 and 316 have two-phase ferrite (3) plus austenite (y) microstructures that are found after primary ferrite solidification under normal welding conditions, these steels often have a fully austenitic microstructure under rapid solidification conditions such as laser-beam welding. This is caused by a change in the solidification mode from the equilibrium primary ferrite solidification mode to a non-equilibrium primary austenite solidification mode due to rapid solidification (Vitek and David 1994). 5. 7.1.2 Non-ferrous alloys The low overall heat input and, thereby, the small fusion and heat affected zones makes laser-beam joining suitable for welding age-hardenable alloys, which have a problem of softening in welded zones, in comparison with conventional arc welding processes. The characteristics of laser-beam welds of age-hardened A1 and Ni base alloys will be hereinafter presented in comparison with their tungsten inert gas (TIG) welds. Age-hardened A1-Mg-Si alloy, 606 l-T6, thin plates (1 mm in thickness) have been welded with full penetration using a 2.5 kW CO2 laser (Hirose et al. 1997b). In laser-beam welding a full penetration weld can be achieved even at a welding speed of 167 mm/s, whereas in TIG welding a welding speed of 5 mm/s or less is required for the full penetration weld. The cross section of a typical laser-beam weld is shown in Fig. 5.9. The bead width of the laser-beam weld is approximately 1/3 that of the TIG weld. Figure 5.10 shows the hardness profiles obtained from the laser-beam and TIG welds in the as-welded condition and after an artificial aging treatment at 448 K for 28.8 ks. In the as-welded condition, although a softened region is found in both the laser-beam weld and the TIG weld, the width of the softened region in the laser-beam weld is less than 1/4 that of the TIG weld. The softened region is caused by reversion (dissolution) of strengthening/~
Laser Processing .
109
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(Mg2Si) precipitates due to the heat input of welding. The hardness of the softened region in the laser-beam weld is almost fully recovered to the level of the base alloy after an aging treatment at 448 K for 28.8 ks. In contrast to this, the TIG weld subjected to the aging treatment after welding has an appreciable hardness drop in a region between 4 and 7 mm from the bead center. This region is thought to be the region where the solute (Mg and Si) content in the matrix is insufficient for precipitation of the strengthening fl (MgzSi) phase because of precipitation of nonhardening fif and/or fl1~ phases during weld thermal cycle. Therefore, the TIG weld will require a solution annealing treatment before artificial aging to fully recover the heat-affected zone hardness unlike the laser-beam welds. In laser-beam welding, since the cooling rate in the heat affected zone is rapid enough to prevent the precipitation of fl~ and/or fl~ phases, supersaturated solid solutions can be obtained over the whole region of the heat affected zone, and thereby sufficient age-hardening due to precipitation of fi1~ phase can be achieved after the artificial aging. As expected from the hardness profile in Fig. 5.10, the tensile strength of the specimen subjected to the aging treatment after laser-beam welding is almost equivalent to that of the base plate. Laser-beam welding has been applied to joining an age-hardened Ni base superalloy, Inconel 718, plate (3.2 mm in thickness) using a 2.5 kW CO2 laser (Hirose et al. 1997a, Hirose et al. 1998) and its properties have been compared with those of TIG welds. Inconel 718 is strengthened by intermetallic ~,~ and V ~ phases, which are precipitated after solution annealing treatment followed by a two step aging treatment. Microstructures of the weld fusion zones are shown in Fig. 5.11. The TIG weld fusion zone has a cellular dendritic structure having austenite (Vf)/Nb-rich Laves phase eutectic and Nb-rich MCtype carbides in the interdendritic regions (Fig. 5.11 (a)). The laser-beam weld fusion zone
110
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has significantly finer solidification structure than that of the TIG weld (Fig. 5.1 l(b)). In the laser-beam welding, the cooling rate during solidification that is estimated from the relationship between dendrite arm spacing and cooling rate is 100 times greater than that in the TIG welding. Thus, the high energy concentration of the laser-beam welding results in rapid solidification and thereby fine solidification structure. Figure 5.12 shows the hardness distributions obtained from the laser-beam and TIG welds in the as-welded condition and after the two-step aging treatment. The fusion zone width of the laser-beam weld is approximately 1/6 of that of the TIG weld. Apparent softening occurs in the weld fusion and heat-affected zones because of a reversion (dissolution) of strengthening y' and y" phases. However, the laser-beam weld has a significantly narrower softened region that is approximately 1/7 of that of the TIG weld. In the laser-beam welds, the hardness of the fusion and heat-affected zones is fully recovered to that of the aged base alloy after the aging treatment. In the TIG welds, the hardness of the weld fusion zones after the aging treatment is not completely recovered to the level of the base metal. Since in the TIG weld fusion zones coarse Nb-rich phases are formed in interdendritic regions, the solute concentrations of the y' matrix may be insufficient for full precipitation of the y' and y" phases. These hardness distributions apparently reveal an advantage of laser-beam welding that its low overall heat input results in fine solidification structure and minimize the weld fusion and heat affected zones. The tensile strength of the laser-beam weld joint subjected to the aging treatment after welding fully recovers to that of the aged base alloy, whereas that of the TIG weld joint does not because of the coarse interdendritic phases and lower hardness in the weld fusion zone.
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111
Figure 5.11. Microstructures of weld fusion zones of Ni-base superalloy, Incone1718. (a) TIG welding and (B) CO2 laser-beam welding.
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112
K. E Kobayashi
Narrow weld-fusion and heat-affected zones, and minimizing distortion in laser-beam welding, are also beneficial to fusion welding, of brittle materials such as intermetallic compounds. Laser-beam welding of a molybdenum-containing TiA1 intermetallic compound (Ti-46at%A1-2at%Mo) has been performed using a 2.5 kW CO2 laser (Hirose et al. 1995a, Hirose et al. 1995b). The cooling rates in the welded zones have been controlled by changing the welding speed and pre-heating temperature. The base metal has the lamellar structure containing y(TiA1) + otz(Ti3A1). However, the microstructure and hardness of the laser fusion zones change depending upon the cooling rates as shown in the schematic CCT diagram (Fig. 5.13). When the cooling rate at 1390 K is above 3000 K/s, the ot ~ y transformation is suppressed and the fusion zone structure results in a single phase of ot2. If the cooling rate is less than 3000 K/s, massive y phase is seen together with the Ot2 phase. In laser-beam welding, where the cooling rates are approximately 80 K/s to 1800 K/s, the fusion-zone microstructure mainly consists of massive ot2, massive y and lamellar (c~2 + y) phases. Thus metastable phases of massive Ot2 and massive y form owing to rapid cooling during laser-beam welding. A fully lamellar structure is formed if the cooling rate is less than 20 K/s. The hardness of the fusion zones increases from 315 Hv of the base metal hardness to above 500 Hv with increasing cooling rate. Cracking in the weld fusion zones is sensitive to the cooling rates and hence to the microstructures of the fusion zones. When the hardness value of the fusion zones is reduced below 410 Hv, cracking dramatically decreases. As a result, crack-free welds can be obtained at welding speeds below 50 mm/s and pre-heating temperatures above 673 K (Fig. 5.14). The cross-section of a crack-free weld bead is shown in Fig. 5.15. During tensile testing, the laser welded specimens without cracking fractured in the base metal. Thus, laser-beam welding with preheating can provide narrow weld fusion and heat-affected zones, and optimal microstructures of the weld fusion zones that are beneficial to welding of brittle materials.
5.7.2 Microjoining
Microjoining in electronics industries also is an important application of laser-beam processing. Laser-beam soldering has been applied to surface-mount technology, namely interconnection between electronic components and printed circuits since 1980's (Chang 1986, Semerad et al. 1993, Messier and Millard 1994). In laser-beam soldering, ease of controlling the shape and location of the heating region brings about reliable solder joints with minimal component heating and is suitable for fine-pitch soldering and high-density packaging. While CO2 laser and Nd:YAG laser are used for laser soldering, Nd: YAG laser is more suitable because of preferential absorption by solders as opposed to board materials (Messier and Millard 1994). Since only the solder joint area is heated during laser-beam soldering process without damaging the boards and components, this process is more beneficial to interconnections using solders with a wide range of melting temperatures in comparison with conventional infrared reflow process (Messier and Millard 1994). Thus, this can overcome the problem of soldering with Pb-free solders with higher melting temperatures than those of eutectic or near eutectic Sn-Pb solders (Yang et al. 1995). Rapid heating and cooling thermal
Laser Processing
113
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114
K. E Kobayashi
Figure 5.15. Cross-sectionof a crack free laser-beamweld of 2 mm thick Ti-46at%A1-
2at%Mo using 2.5 kW CO2 laser.
cycles in laser-beam soldering result in the microstructures and mechanical properties of the solder joints different from those of the conventional infrared reflowed joints. Laser-beam soldering of eutectic Sn-Ag solder has been performed and the effect of the soldering process on the microstructures and mechanical properties have been evaluated (Yang et al. 1995). In this study eutectic Sn-Ag solder joints between quad-flat pack surface-mounted components and printed wiring boards have been produced using a Nd : YAG laser soldering system. Much lower overall heat input and shorter duration time in laser-beam soldering results in very high cooling rates (~104 K/s) from peak temperature compared with the infrared reflow soldering (~10 K/s). Therefore, in laserbeam solder joints, much finer eutectic Ag3Sn intermetallic in pure tin and thinner Cu-Sn intermetallic layers at copper pad and wire interfaces form compared with the infrared
Laser Processing
115
reflowed joints. The high peak temperature (673 K to 873 K) caused by laser irradiation compared to the infrared reflow (~528 K) dissolves a certain amount of copper from leads and pads. This results in the formation of a considerable number of equiaxed Cu6Sn5 dendrites in the bulk solder. The difference of microstructure between laser-beam soldered and infrared reflowed joints causes higher microhardness of the laser-beam soldered joints (~27 Hv) than that of infrared reflowed joints (~16 Hv), suggesting better mechanical properties under static, creep and fatigue loading. Laser-beam processes have also been applied to bond Tape Automated Bonding (TAB) leads to integrated circuit pads (Spletter and Crowley 1990). Since laser inner lead bonding can be performed using lower temperature and lower pressure compared to conventional gang bonding process, it can minimize the thermal and mechanical stressinduced damages of circuit. TAB laser bonding is performed by positioning a focused laser-beam over a lead and then pulsing the laser once, creating a reliable metallurgical bonding between the plated lead and the bump (Spletter and Crowley 1990). This bonding process takes advantage of a thin tin plating to absorb the laser energy and accelerates melting of the copper lead and gold bump. Since this process achieves rapid alloying of copper and gold, bonds without the brittle intermetallic compounds often found in gold/tin alloys formed with conventional methods can be obtained (Spletter and Crowley 1990).
5.8. C O N C L U S I O N S Lasers are lights that obey straightforward physical laws and make it easy to predict their behavior. Knowing the relevant optical equations and how to apply them will explain most of the features of laser materials processing. For example, the high energy density available from focused sources is utilized to melt only the surface of a material in a very short time. This characteristic can apply to form non-equilibrium materials. In other cases like EXCIMER laser, it is suitable for applications in the field of micromachining. Although laser processing is prominent, the user needs to consider the following problems. What is the best quality that one can achieve and what is the fastest way to get there? What will be new and different are the orders of magnitude differences between the feature sizes and quality. These differences will demand better optics, stable environments, and system design, part handling and inspection systems.
REFERENCES
Akgun, O. V. and Inal, O. T. (1995) J. Mater. Sci., 30, 6097. Aziz, M. J. (1982) J. AppL Phys., 53, 1158. Badan, B., Giordano, L., Ramous, E. and Rudilosso, S. (1985) in Rapidly Quenched Metals V, eds. Steeb, S. and Warlimont, H. (Elsevier Science Pub., Amsterdam),p. 851. Bergmann, H. W. (1985) Surface Eng., 1, 135. Bergmann, H. W. and Mordike, B. L. (1981) J. Mater Sci., 16, 863. Bergmann, H. W., Fritsch, H. U. and Hunger, G. (1981) J. Mater. Sci., 16, 1935. Bimberg, D. (1986) in Laser Treatment of Materials, ed. Mordike, B. L. (DGM Informationsgesellschaft, Bad Nauheim, Germany),p. 3. Bloembergen, N. (1979) in Laser-Solid Interactions and Laser Processing 1978, eds. Ferris, S. D., Leamy, H. J. and Poate, J. M. (Amer. Inst. Phys., New York), p. 1.
116
K. F. K o b a y a s h i
Boettinger, W. J., Biancaniello, E, Kalonji G. M. and Cahn, J. W. (1980) in Rapid Solidification Processing: Principles and Technologies, eds. Mehrabian, R., Kear, B. H. and Cohen, M. (Claitor's Pub., Reston, VA), p. 50. Boettinger, W. J., Coriell, S. R. and Sekerka, R. E (1984) Mater. Sci. Eng., 65, 27. Boettinger, W. J., Coriell, S. R. and Trivedi, R. (1988) in Rapid Solidification Processing: Principles and Technologies, eds. Mehrabian, R. and Paris, P. A. (Claitor's Pub., Reston), p. 13. Bradley, J. R. and Kim, S. (1988) Metall. Trans. A, 19A, 2013. Cahn, J. W. (1956) Acta Metall., 4, 572. Carrad, M., Gremaud, M., Zimmermann, M. and Kurz, W. (1992) Acta Metall. Mater., 40, 983. Chande, T. and Mazumder, J. (1982) Appl. Phys. Lett., 41, 42. Chande, T. and Mazumder, J. (1983) Metall. Trans., 14B, 181. Chang, D. U. (1986) Weld. J., 65, 32. Cline, H. E. and Anthony, T. R. (1977) J. Appl. Phys., 48, 3895. Cullis, A. G., Webber, H. C. and Chew, N. G. (1982) in Laser and Electron-Beam Interraction with Solids, eds. Appleton, B. R. and Celler, G. K. (North-Holland, Amsterdam), p. 131. David, S. A. and Vitek, J. M. (1985) in Rapidly Quenched Metals V, eds. Steeb, S. and Warlimont, H. (Elsevier Science Pub., Amsterdam), p. 847. David, S. A. and Vitek, J. (1989) Internat. Mater. Rev., 34, 23. David, S. A., Vitek, J. M. and Hebble, T. L. (1987) Weld. J., 68, 289-s. Davies, H. A., Aucote, J. and Hull, J. B. (1974) Scripta Metall., 8, 1179. Draper, C. W. and Poate, J. M. (1985) lnternat. Metall. Rev., 30, 85. Egami, T. and Waseda, Y. (1984) J. Non-Cryst. Solids, 64, 113. Elhossary, E M., Fabian, D. J. and Webb, A. P. (1990) Thin Solid Films, 192, 201. Frauenheim, E M., Stephan, U., Bewilogua, K., Jungnivkel, E, Blaudeck, P. and Fromm, E. (1989) Thin Solid Films, 182, 63. Fritsch, H. U. (1986) in Laser Treatment of Materials, ed. Mordike, B. L. (DGM Informationsgesellschaft, Bad Nauheim), p. 19. Garcia-Alonso, M. C., Escudero, M. L., Lopez, V. and Macias, A. (1996) ISIJ lnternat., 36, 172. Gilgien, P., Zryd, A. and Kurz, W. (1995) Acta Metall. Mater., 43, 3477. Gill, S. C. and Kurz, W. (1995) Acta Metall. Mater., 43, 139. Grange, R. A. and Kiefer, J. M. (1941) Trans. ASM, 29, 85. Greenwald, L. E., Breinan, E. M. and Kear, B. H. (1979) in Laser-Solid Interactions and Laser Processing 1978, eds. Ferris, S. D., Leamy, H. J. and Poate, J. M. (Amer. Inst. Phys., New York), p. 189. Herzer, G. (1992) J. Magn. Magn. Mater., 112, 258. Hirose, A. and Kobayashi, K. E (1994) Mater Sci. Eng., A174, 199. Hirose, A. and Kobayashi, K. F. (1995) ISIJ Internat., 35, 757. Hirose, A., Ueda, T. and Kobayashi, K. E (1993) Mater. Sci. Eng.,A160, 143. Hirose, A., Arita, Y. and Kobayashi, K. E (1995a) J. Mater Sci., 30, 970. Hirose, A., Arita, Y. and Kobayashi, K. E (1995b) J. Society Mater Sci. Japan, 44, 1145. Hirose, A., Arita, Y., Nakanishi, Y. and Kobayashi, K. E (1996) Mater Sci. Eng., A219, 71. Hirose, A., Sakata, K. and Kobayashi, K. E (1997a) in Proc. of 4th Decennial Int. Conf. on Solidification Processing, Sheffield, UK, p. 675. Hirose, A., Todaka, H. and Kobayashi, K. E (1997b) Metall. Mater. Trans., 28A, 2657. Hirose, A., Sakata, K. and Kobayashi, K. E (1998) Int. J. Mater. Production Technol. (submitted). Hirose, E, Takagi, M., Moil, H. and Imura, T. (1992) Jpn. J. Appl. Phys., 31, 3940. Hirose, E, Takagi, M., Mori, H. and Imura, T. (1994) Mater Sci. Eng., A181/A182, 1359. Hsu, S. C., Chakravorty, S. and Mehrabian, R. (1978) Metall. Trans., 9B, 221. Hsu, S. C., Kou, S. and Mehrabian, R. (1980) Metall. Trans., liB, 29. Hunziker, O. and Kurz, W. (1997) in Proc. of 4th Decennial Internat. Conf. on Solidification Processing, Sheffield, UK, p. 599. Ion, J. C., Salminen A. S. and Sun, Z. (1996) Weld. J., 75, 225-s. Jackson, K. A., Gilmer, G. H. and Leamy, H. J. (1980) in Laser and Electron Beam Processing of Materials, eds. White, C. W. and Peercy, P. S. (Academic Press, New York), p. 104. Khan, P. A. A., Debroy, T. and David, S. A. (1988) Weld. J., 67, 1-s. Kobayashi, K. E and Shingu, P. H. (1982) in Rapidly Quenched Metals IV, vol. I, eds. Masumoto, T. and Suzuki, K. (Japan Inst. Metals, Sendai), p.103.
Laser Processing
117
Kobayashi, K. E and Shingu, E H. (1988) J. Mater. Sci., 23, 2157. Kobayashi, K. E, Kumikawa, M. and Shingu, E H. (1984) J. Cryst. Growth, 67, 85. Kou, S., Hsu, S. C. and Mehrabian, R. (1981) Metall. Trans., 12B, 33. Kumagai, N., Asami, K. and Hashimoto, K. (1986) J. Non-Cryst. Solids, 87, 123. Kuriyama, H., Kiyama, S., Noguchi, S., Kuwahara, T., Ishida, S., Nohda, T., Sano, K., Iwata, H., Kawata, H., Osumi, M., Tsuda, S., Nakano, S. and Kuwano, Y. (1991) Jpn. J. Appl. Phys., 30, 3700. Kuriyama, H., Kuwahara, T., Ishida, S., Nohda, T., Sano, K., Iwata, H., Noguchi, S., Kiyama, S., Tsuda, S., Nakano, S., Osumi, M. and Kuwano, Y. (1992) Jpn. J. Appl. Phys., 31, 4550. Kuriyama, H., Nohda, T., Ishida, S., Kuwahara, T., Noguchi, S., Kiyama, S., Tsuda, S. and Nakano, S. (1993) Jpn. J. Appl. Phys., 32, 6190. Kurokawa, H., Yonezawa, T. and Mitani, T. (1987) IEEE Trans. Magnetics, 23, 2410. Kurz, W. (1995) in Mathematical Modelling of Weld Phenomena, ed. Cerjak, H. (The Institute of Metals, London), p. 40. Kurz, W., Giovanola, B. and Trivedi, R. (1986) Acta Metall., 34, 823. Kurz, W. and Trivedi, R. (1994) Mater Sci. Eng., A179/180, 46. Lanotte, L. and Lannotti, V. (1995) J. Appl. Phys., 78, 3531. Lee, H. J., Akiyama, E., Habazaki, H., Kawashima, A., Asami, K. and Hashimoto, K. (1995) Corros. Sci., 37, 321. Lin, C. J. and Spaepen, E (1981) Appl. Phys. Lett., 41, 55. Lin, C. J. and Spaepen, E (1982) Appl. Phys. Lett., 41, 721. Lin, C. J. and Spaepen, E (1983) in Chemistry and Physics of Rapidly Solidified Materials, eds. Berkowitz, B. J. and Scattergood, R. (TMS-AIME, New York), p. 273. Lin, C. J. and Spaepen, E (1984) Mater Res. Soc. Symp. Proc., 28, 75. Lin, C. J. and Spaepen, E (1986) Acta Metall., 34, 1367. Lin, C. J., Spaepen, E and Turnbull, D. (1984) J. Non-Cryst. Solids, 33, 251. Lipton, J., Kurz, W. and Trivedi, R. (1987) Acta Metall., 35, 957. Liu, A. Y. and Cohen, M. L. (1990) Phys. Rev., B41, 10727. Liu, E L., Yen, R., Bloembergen, N. and Hodgson, R. T. (1979) Appl. Phys. Lett., 34, 864. Lowndes, D. H., Pennycook, S. J., Jellison, Jr., G. E., Withrow, S. E and Washburn, D. N. (1987) J. Mater. Res., 2, 648. Lu, E X., Yang, B. X., Cheng, D. G., Ye, R. Z., Yu, W. X. and Sun, J. B. (1992) Thin Solid Films, 212, 220. Matteazzi, E, Lanotte, L. and Tagliaferri, V. (1985) in Rapidly Quenched Metals V, eds. Steeb, S. and Warlimont, H. (Elsevier Science Pub., Amsterdam), p. 1199. Mazumder, J. and Steer, W. M. (1980) J. Appl. Phys., 51, 941. Mazzoldi, E, Della Mea, G., Battaglin, G., Miotello, A., Servidori, M., Bacci, D. and Jannitti, E. (1980) Phys. Rev. Lett. 44, 88. Mehrabian, R. (1982) Internat. Metall. Rev., 27, 185. Mehrabian, R., Kou, S., Hsu, S. C. and Munitz, A. (1979) in Laser-Solid Interactions and Laser Processing 1978, eds. Ferris, S. D., Leamy, H. J., and Poate, J. M. (Amer. Inst. Phys., New York), p. 129. Messler, Jr., R. W. and Millard, D. L. (1994) Weld. J., 73, 43. Metzbower, E. A. (1990) Crystal Properties and Preparation, 22, 656. Metzbower, E. A. and Denney, E E. (1986) Proc. Inernat. Conf. Power Beam Technol., p. 10-1. Mitsuhashi, A., Asami, K., Kawashima, A. and Hashimoto, K. (1987) Corros. Sci., 27, 957. Miyazawa, Y. and Pound, G. M. (1974) J. Cryst. Growth, 23, 45. Moon, D. W., Metzbower, E. A., Phillips, R. H. and Denney, E E. (1987) in Proc. Conf. Laser vs. Electron Beam in Welding, Cutting and Surface Treatment: State of the Art, Part 2, p. 280. Pound, G. M. and Lamer, V. K. (1974) J. Amer. Chem. Soc., 74, 2323. Prasad, S. V., Zabinski, J. S. and Dyhouse, V. J. (1992) J. Mater. Sci. Lett. 11, 1282. Rappaz, M., Carrupt, B., Zimmermann, M. and Kurz, W. (1987) Helvetica Physica Acta, 60, 924. Renaud, L., Fouquet, E, Millet, J. E, Mazille, H. and Crolet, J. L. (1991) Mater Sci. Eng., A134, 1049. Roessler, D. M. (1986) in The Industrial Laser Handbook, eds. Belforte, D. and Levitt, M. (Penny Well Books, Tulsa), p. 16. Rosenthal, D. (1946) Trans. ASME, 68, 849. Rozgonyi, G. A., Baumgart, H., Phillip, E, Uebbing, R. and Oppholzer, H. (1982) in Laser and Electron-Beam Interraction with Solids, eds. Appleton, B. R. and Celler, G. K. (North-Holland, Amsterdam), p. 177. Sadler, L. Y. and Shamsuzzoha, M. (1997) J. Mater Res., 12, 147.
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Sameshima, T., Usui, S. and Sekiya, M. (1986) IEEE Electron Device Lett., 7, 276. Sato, T., Yamada, T. and Ozawa, T. (1985) in Rapidly Quenched Metals V, eds. Steeb, S. and Warlimont, H. (Elsevier Science Pub., Amsterdam), p. 1643. Scandella, L., Staufer, U., Brodbeck, D., Reimann, P., Gtintherodt, H. J., Zehringer, R., Hauert, R. and Moser, E. M. (1991) Mater Sci. Eng., A133, 601. Schaeffler, A. L. (1949) Metal Prog., 56, 680. Schwan, J., Dworschak, W., Jung, K. and Ehrhardt, H. (1994) Diamond Relat. Mater., 3, 1034. Sekhar, J. A. and Mehrabian, R. (1981) Metall. Trans., 12B, 411. Semerad, E., Musiejovsky, L. and Nicolas, J. (1993) J. Mater. Sci., 28, 5065. Sera, K., Okumura, E, Uchida, H., Itoh, S., Kaneko, S. and Hotta, K. (1989) IEEE Trans. Electron Devices, 36, 2868. Shimizu, K., Sugiura, O. and Matsumura, M. (1993) IEEE Trans. Electron Devices, 40, 112. Singh, J. and Mazumder, J. (1987a) Metall. Trans., 18A, 1987. Singh, J. and Mazumder, J. (1987b) Acta Metall., 35, 1995. Spaepen, E (1975) Acta Metall., 23, 729. Spaepen, E and Turnbull, D. (1976) in Rapidly Quenched Metals, eds. Grant, N.J. and Giessen, B. C. (MIT Press, Cambridge), p. 205. Spaepen, E and Turnbull, D. (1982) in Laser Annealing of Semiconductors, eds. Poate, J. M. and Mayer, J. W. (Academic Press, New York), p. 15. Spletter, P. J. and Crowley, R. T. (1990) in Proc. Electronics Components & Technology Conf., p. 757. Tabata, H. and Kawai, T. (1997a) Appl. Phys. Lett., 70, 321. Tabata, H. and Kawai, T. (1997b) IEICE Trans. Electron, E-80-C, 918. Tabata, H., Kawai, T. and Kawai, S. (1993) Phys. Rev. Lett., 70, 2633. Tabata, H., Tanaka, H. and Kawai, T. (1994) Appl. Phys. Lett., 65, 1970. Tabata, H., Tanaka, H. and Kawai, T. (1995) Jpn. J. Appl. Phys., 34, 5146. Trivedi, R., Magnin, P. and Kurz, W. (1987) Acta Metall., 35, 971. Tsu, R., Hodgson, R. T., Tan, T. Y. and Baglin, J. E. (1979) Phys. Rev. Lett., 42, 1356. Tsukamoto, H., Yamamoto, H., Noguchi, T. and Suzuki, T. (1992) Jpn. J. Appl. Phys., 31, L659. Tsukamoto, H., Yamamoto, H., Noguchi, T., Masuya, H. and Suzuki, T. (1996) Jpn. J. Appl. Phys., 35, 3810. Turnbull, D. (1950) J. Appl. Phys., 21, 1022. Uenishi, K. and Kobayashi, K. E (1993) in Proc. 3rd Japan Internat. SAMPE Symposium, p. 1038. Uenishi, K. and Kobayashi, K. E (1994) in Proc. of the 4th lnternat. Conf. on Aluminum Alloys, (Atlanta, GA), p. 714. Uenishi, K. and Kobayashi, K. E (1996a) Solidification Science and Processing, eds. Ohnaka, I. and Stefanescu, D. M. (TMS, Warrendale), p. 231. Uenishi, K. and Kobayashi, K. E (1996b) Intermetallics, 4, $95. Uenishi, K., Sugimoto, A. and Kobayashi, K. E (1992a) Z. Metallkde., 83, 241. Uenishi, K., Sugimoto, A. and Murakami, T. (1992b) in Proc. Int. Conf. on Laser Advanced Materials Processing (Nagaoka, Japan), p. 807. Uhlman, D. R. (1972) J. Non-Cryst. Solids, 7, 337. Vitek, J. M. and David, S. A. (1994) Laser Materials Processing-IV, eds. Mazumder, J., Mukherjee, K. and Mordike, B. L. (TMS, Warrendale), p. 153. Vitek, J. M., Das Gupta, A. and David, S. A. (1982) Metall. Trans. 13A, 1141. Vitta, S., Greer, A. L. and Somekh, R. E. (1994) Mater. Sci. Eng., A179/A180, 243. von Allmen, M. (1983) in Glassy Metals 1I, eds. Beck, H. and Gtintherodt, H.-J. (Springer-Verlag, Berlin), p. 261. von Allmen, M., Lau, S. S., Maenpaa, M. and Tsaur, B. Y. (1980) Appl. Phys. Lett., 37, 84. Yang, W., Messier, Jr., R. W. and Felton, L. E. (1995) Weld. J. 74, 224-s. Yano, T., Ooie, T., Yoneda, M. and Katsumura, M. (1996) J. Mater. Sci. Lett. 15, 1994. Yee, A. L., Ong, H. C., Xiong, F. and Chang, R. P. H. (1996) J. Mater. Res. 11, 1979. Young, R. T. and Wood, R. E (1982) Ann. Rev. Mater. Sci., 12, 323. Young, R. T., Leeden, G. A., Narayan, J., Christie, W. H., Wood, R. E, Rothe, D. E. and Levatter, J. I. (1982) IEEE Electron Device Lett., 3, 280. Zabinski, J. S., Donley, M. S., Prasad, S. V. and McDevitt, N. T. (1994) J. Mater. Sci. 29, 4834. Zambon, A. and Bonollo, E (1994) Mater. Sci. Eng., A178, 203. Zarubova, N., Wolf, P., Cermak, J. and Cernansky, M. (1996) J. Mater. Sci. 31, 137.
Chapter 6 Thermal Plasma Processing 6.1.
Introduction 6.1.1 Advantages of Plasma Processing 6.2. Thermal Plasmas 6.2.1 Principles of Plasma Generation 6.2.1.1 DC Plasma Torches 6.2.1.2 AC Plasma Torches 6.2.1.3 RF Plasma Torches 6.2.2 Plasmagen Gases 6.2.3 Plasma-Particle Interaction 6.2.4 Plasma Processing Systems 6.3. Processing of Materials 6.3.1 Plasma-Spraying 6.3.1.1 Structure of Sprayed Deposits 6.3.1.2 Plasma Spray-Deposited Materials 6.3.1.3 Low-Pressure Plasma-Spraying 6.3.1.4 Reactive Plasma-Spray Forming 6.3.1.5 Plasma Spheroidization 6.3.2 Plasma Reactors and Furnaces 6.3.2.1 Plasma Decomposition 6.3.2.2 Plasma Metallurgy 6.3.2.3 Processing of Ceramics 6.3.2.4 Treatment of Hazardous Wastes 6.3.3 Processing of Metastable Phases 6.3.3.1 Plasma Deposition of Diamonds 6.3.3.2 Thermal Plasma Synthesis of Ultrafine Alumina Powder 6.4. Summary and Conclusions Acknowledgments References
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Chapter 6 Thermal Plasma Processing R V. ANANTHAPADMANABHAN and N. VENKATRAMANI
6.1. INTRODUCTION The twentieth century, variously termed at different times the semiconductor age, the silicon age, and the oxide age, has seen some breathtaking achievements in the development of new materials. These achievements have been made possible by matching developments in the material processing technologies. Today's advanced industries need precisely tailored materials to meet stringent operational requirements. However, a close look at the conventional manufacturing technology tells us that it is beset with barriers, both theoretical and man-made. Table 6.1 summarizes the limitations of the current material processing technologies and how they can be overcome in the emerging technologies of the future. It may be seen that many of the shortcomings of conventional technologies can be overcome by the use of electron, plasma and laser beams. In particular, thermal plasma processing has the capability to transcend thermodynamic and kinetic barriers by opening up new reaction paths inconceivable in the conventional sense. This technology, which until recently was confined to the realms of laboratory, is fast emerging as a major technique not only to synthesize and purify materials, but also to modify and shape surfaces (Boulos 1996, Venkatramani 1995, Kong and Lau 1990, Pauloski 1995). Plasmas of technological interest can be classified into two categories: nonequilibrium plasmas and thermal plasmas (Fauchais et al. 1983, Fineman 1987). Nonequilibrium or cold plasmas, more popularly known as glow-discharge plasmas, are low-pressure plasmas characterized by high electron temperatures and low ion and neutral temperatures. They are widely used in lighting, surface cleaning, etching, film deposition and polymerization. Thermal plasmas are characterized by the electron temperature being equal to the gas temperature. Normally, plasmas in the temperature range of 2,00030,000 K and with charged particle density of 1019-1021 m -3 are termed thermal plasmas. Thermal plasma processing has emerged as a major processing route to synthesize many alloys and ceramics of industrial importance. The different categories of plasmas and their applications are listed in Table 6.2.
6.1.1 Advantages of plasma processing The high temperature, high enthalpy, and fast quench rate which are characteristics of the process, offer unlimited scope to the materials scientist to synthesize newer ma121
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P V. Ananthapadmanabhan and N. Venkatramani
Table 6.1. Limitations of current material processing technologies Barrier
Approach to overcome the barrier
Emerging technology
Thermodynamic barrier
Processing in non-equilibrium and thermal plasmas
Spatial barrier
Access to non-equilibrium regimes, use of exotic intermediates and precursors Transfer heat to volume rather than across surface Micro- and nanoscale processing
Temporal barrier
Nano- and picosecond processing
Purity barrier
Absolute chemical control and limit impurity content
Heat transfer barrier
and
specificity
Ecological barrier
One-step eco-friendly technology, conserving materials and recycling wastes
Electromagnetic processing Beam processing; electron and laser beams and low-pressure plasma processing Laser processing, thermal plasma processing with rapid quench Processing by electron, plasma and laser beams and low-pressure plasmas Thermal plasma processing
Table 6.2. Plasmas of technological interest Category of plasma
Applications
Low pressure plasmas (10-4-10 -2 torr)
Sputtering and surface modification processes, plasma source for ion implantation Etching, microelectronic processing Plasma chemistry, plasma polymerization
Medium pressure plasmas (10-2-1 torr) Subatmospheric pressure plasmas (1-100 torr) Atmospheric plasmas (100+ torr)
Plasma-spraying, plasma-melting, material synthesis
terials with improved properties. The temperature encountered in the plasma medium is 10,000-15,000 K, whereas the maximum temperature available in conventional heat sources is 3,000 K. This enables the adoption of novel approaches to accomplish chemical reactions and metallurgical processes, inconceivable in the conventional sense. The high-temperature gas phase chemistry in the plasma medium promotes various chemical reactions at much faster rates than those encountered in conventional processing techniques. The high quench rate, typically of the order of 105-107 K/s, prevents recombination of the products and favors homogeneous nucleation resulting in nano-sized particles. Furthermore, it offers flexibility in the choice of reactants, which can be solids, liquids or gases. Plasma processes are, usually, one-step processes, with the capability of handling large throughputs in small reactor volumes in relatively short processing times. They also do not substantially contaminate the process environment, thereby ensuring high purity of the product.
Thermal Plasma Processing
123
Plasma processing has been applied to synthesize a variety of alloys, intermetallics, and refractory ceramics. Plasma-spray technology is a well-established surface treatment technique to prepare protective coatings for a variety of industrial applications. The focus of this chapter will be on thermal plasma processing of materials. The organization of the chapter is as follows: Introduction to the subject of thermal plasma processing will be followed by a description of the different techniques of plasma generation. Applications of thermal plasmas in material processing are discussed in detail in Section 6.3. Plasma synthesis of ultrafine alumina is discussed in detail towards the end of this section to illustrate the formation of non-equilibrium phases in reactions occurring in plasma medium.
6.2. T H E R M A L
PLASMAS
The classical plasma is a fully ionized gas consisting of electrons and ions that interact with each other only through Coulomb forces (Eliezer and Eliezer 1989). This model is valid only for astrophysical or fusion plasmas; it does not hold good for industrial plasmas, which are partially ionized gases and are dominated by atomic and molecular processes. Thermal plasma is a viscous, electrically and thermally conducting fluid. Plasma retains many of the properties of gases, and behaves in conformance with the physical laws valid for gases. The specific properties of plasma that distinguish it from a gas become apparent in the presence of a magnetic field, when the plasma acquires non-isotropic properties. Thus, the fundamental difference between plasma and a non-ionized gas lies in their responses to electromagnetic forces. The electrically charged particles present in plasma are affected by externally applied electric and magnetic fields, and also interact with one another. The electric field that they set up is so extensive that every particle is affected by a multitude of other particles. Consequently, the collective behavior of the particles is much stronger in plasma than in a non-ionized gas. The unique feature of the thermal plasma jet that distinguishes it from other heat sources is its high power density. The energy density of thermal plasma devices is of the order of several GW/m 2, which is 10-100 times the power density of conventional oxy-fuel flames. Although higher power densities are obtainable with electron and laser beams, the source strengths available with these devices, especially laser devices, are not high enough for large-scale material processing applications.
6.2.1 Principles of plasma generation Thermal plasmas for material processing applications are usually of electric origin. Laboratory thermal plasmas are produced in devices known as plasma torches or plasmatrons (Heberlein 1992). Plasma torches are electrothermal devices, which convert the electrical input energy to thermal energy of the plasma. Depending on the primary source, which can be direct current (DC), alternating current (AC) at mains frequency, or at radio frequency (RF), they are known as DC, AC or RF torches. There are also combinations of these torches known as hybrid devices. Figure 6.1 gives a schematic of the different types of plasma torches.
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6.2.1.1 DC plasma torches The most commonly used plasma device for material processing applications is the DC arc plasma torch. A direct current arc is struck between a cathode rod and a nozzle anode and the energy from the arc is abstracted by the plasma gas injected in the inter-electrode region. Forced gas flow extends the arc in the anode nozzle which is strongly water-cooled (Fig. 6.1(a)). A thermal arc pinch effect is produced by the joint action of the cold wall arc channel and the cold gas sheath around a very high-temperature conducting core (the arc column). The DC arc in a plasma torch needs to be stabilized, i.e., it should remain stationary against fluctuations. This is often done by constricting the arc to a well-adjusted narrow high-temperature, highly conducting arc column. Various torch configurations are possible depending upon the arc stabilization mode: tangential vortex gas input in the arc channel, axial gas input along the cathode, segmented anode arc and magnetic stabilization (Fauchais et al. 1983). The magnetic field can be self-induced (by an arc current greater than 8000 A), or externally generated. The electric arc is stabilized by using a constricted anode nozzle and by the resultant aerodynamic effects in the streaming plasma gas. Stabilizing action of the vortex gas flow provides a cold boundary layer near the anode wall so that heat loss to the wall is reduced. This results in the thermal energy being highly concentrated with improved torch stability and efficiency. The localized heat flux at the nozzle can be as high as 160 W/mm 2. The obvious choice of material for the anode is copper, although molybdenum and graphite are also used. Depending on the nature of the gas and the working parameters, the anode losses range
Thermal Plasma Processing 3
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between 80 and 10% of the energy input in the arc. The anode losses are proportional to the current density and are a function of the arc voltage. The cathodes can be of the thermionic type such as tungsten, carbon or molybdenum, which, obviously, must be used in non-oxidizing atmosphere. Under certain conditions of oxidizing atmosphere, one can use zirconium or hafnium cathodes. Heat losses at the cathode are generally quite low (less than 10% of the input power). For a tungsten cathode tip, the erosion rate has been estimated to be of the order of 10 -9 to 10-10 g per Coulomb, so that its lifetime is about 150 h. Non-thermionic cathodes are normally made of copper. The cross-section of a typical DC plasma torch is shown in Fig. 6.2. This design is based on a rod type cathode and nozzle-shaped anode with tangential gas entry through the insulator module. When a gas is injected into the electrode gap and a high intensity current is passed, a DC arc is established between the electrodes. The plasma gas extracts energy from the arc and emerges out of the nozzle as a high-temperature, high velocity jet. The temperature at the core of the plasma jet ranges between 15,000-20,000 K (Gerdeman and Hecht 1972). The typical operating parameters of a low power, laboratory DC plasma torch (non-transferred arc mode) are given in Table 6.3 The plasma torch can be operated in a transferred or non-transferred mode, depending on whether the arc is electrically transferred to the work piece or not, as illustrated in Fig. 6.1 (a) and 6.1 (b). The transferred arc mode is mainly used for cutting metal sheets, refining, and remelting operations. The non-transferred arc mode is used in plasma-spraying and other material processing operations.
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P. V. Ananthapadmanabhan and N. Venkatramani
Table 6.3 Typical operating parameters of a low-power laboratory DC plasma torch (non-transferred arc mode) Arc voltage (V) Arc current (A) Arc gas Torch input power (kW) Water Inlet temperature ((C) Water outlet temperature ((C) Cooling water flow rate (1/min) Power loss at the electrodes (kW) Flow rate of Argon (g/s) Cold gas flow velocity (m/s) Thermal energy of emerging plasma (kW) Enthalpy of plasma (kJ/g) Plasma temperature at the nozzle exit (K) Plasma exit velocity (m/s)
40 250 Argon 10 33.2 36.3 15 3.2 1.19 13.27 6.8 5.71 10,600" 468.9*
* Calculated.
6.2.1.2 AC plasma torches When alternating current at mains frequency is used in a two-electrode device, each electrode alternately acts as the cathode and anode. In each half cycle a DC discharge is established once the breakdown potential is crossed and the discharge is maintained so long as the applied voltage remains above the maintenance potential. When the voltage falls below the maintenance potential, the discharge gets extinguished, and the space charge gets swept away. The continuous alternation of the polarity of the electrodes limits strongly the stability of the discharge and the choice of the electrode materials.
6.2.1.3 RF plasma torches Radio-frequency (RF) inductive or capacitive coupling (between 1-20 MHz) can be used to generate a plasma column, which can be stabilized in a quartz tube. The main advantage of this technique is its electrodeless feature, allowing the use of aggressive gases such as oxygen, chlorine (Mackinnon and Reuben 1975) or UF6, etc. which cannot be used as plasmagen gases in DC plasma torches. The minimum power necessary to sustain the discharge is determined by the nature of the gas, its pressure and the frequency of the electromagnetic field (Rykalin 1978). The gas velocity in an RF torch is several tens of m/s, which is much lower than that encountered in DC torches (hundreds of m/s) (Dembovski 1985). Furthermore, the diameter of the reactor is important (20-200 mm) and the energy density is much lower than in arc plasma generators. RF plasma torches capable of operating up to power levels of 1 MW (Ionarc Tafa torches) and a lifetime of a few thousand hours are presently employed for material processing applications. However, the electrothermal efficiency of these torches is low (40-50%) and are mostly limited to specialized applications.
Thermal Plasma Processing
,•,4
11,S00K
10,500K
9,600 K
8,600 K
7,600 K
127
6,800 K
E2
0
4
8
12
16
20
24
Axial Distance(mm) Figure 6.3. Axial and radial temperature distribution in a nitrogen plasma jet (Fauchais et al. 1983).
The arc column can be extended by using hybrid torches, which consist of a combination of torches connected in tandem. The extended hot zone increases the residence time and ensures that the reaction goes to completion.
6.2.2 Plasmagen gases The most commonly used gases for plasma generation are argon, nitrogen, helium, hydrogen, and air. Plasma gas flow rate and the electric power to the plasma torch must be properly balanced in order to get a stable arc. The choice of plasma gas depends on many factors, such as the design features of the torch, in particular the electrode materials (Venkatramani 1995). In the case of plasma torches employing tungsten cathode, the choice of the plasma gas is limited to inert gases and non-oxidizing gases. Gas enthalpy is another important factor deciding the choice of the gas (Gerdeman and Hecht 1972). The energy content of nitrogen and hydrogen, which are diatomic, is considerably higher than that of argon or helium. This is due to the dissociation reaction in the case of nitrogen and hydrogen prior to ionization. The heat transfer coefficient is also higher, enabling efficient plasma-particle heat transfer. The low cost and high internal energy of nitrogen makes it the most commonly used gas. If a completely inert atmosphere is required, argon is usually preferred. Reactive gases like hydrogen, oxygen, chlorine, methane and ammonia/nitrogen can be used to impart reducing, oxidizing, chloriding, carburizing, or nitriding effects, respectively, to the plasma. The axial and radial temperature fields in the plasma jet issuing out of a nontransferred arc torch are shown in Fig. 6.3. It is seen from the figure that the temperature falls off rapidly downstream the axis and radially away from the centre-line. The maximum temperature and the temperature profile depend on the input power, nature and flow rate of the plasma gas and the torch design. But the significant point to be noted is the steep gradient in the axial and radial temperatures. Particle injection into the plasma stream is not a trivial problem. In fact, it can be very difficult due to the high viscosity of the plasma. Precursor powders are usually entrained
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P V. Ananthapadmanabhan and N. Venkatramani
in a carrier gas for injection into the plasma. The carrier gas can be inert or reactive. There is an increasing tendency to use liquid reactants, which is known as liquid injection plasma synthesis (LIPS) (Kong and Lau 1990).
6.2.3 Plasma-particle interaction The heat transfer process from the plasma to the particles is of fundamental importance in all plasma processing applications. Plasma-particle heat transfer is very complicated and depends on the thermo-physical properties of the material as well as the enthalpy and composition of the plasma. Many theoretical models (Fauchais et al. 1983, Bonet 1976) have been developed for the plasma-particle heat transfer phenomenon. The heattransfer process is decided by the thermal boundary at the particle-plasma interface and the properties of the gases in this boundary layer, which is characterized by steep temperature gradient and strong deviations from chemical and kinetic equilibrium (Pfender 1994). The heat transfer coefficient to the particle may be calculated from a correlation based on the well-known Ranz-Marshall equation (Ranz and Marshall 1952): Nu = 2.0 + 0.6Re~
~
(6.1)
where Nu, Re and Pr are the Nusselt, Reynolds and Prandtl numbers, respectively. Although this equation must be corrected by additional factors to account for the strong variation of gas properties in the boundary layer, it does give a fair estimate of the heat transfer coefficient. The coefficient of heat transfer between the plasma gas and the surface of the particle can now be calculated from the definition of the Nusselt number: Nu - hdp/k
(6.2)
where h is the heat transfer coefficient, k is the thermal conductivity of the plasma, and
dp is the diameter of the particle. The goal pursued while optimizing the process parameters is to ensure that all particles entering the plasma jet are heated to the desired temperature, which may be the melting point in the case of plasma-spraying or the reaction temperature in chemical synthesis or dissociation reactions. If the particles are small or if they conduct heat well, this objective is easier to achieve. On the other hand, in the case of large or poorly conducting particles, it might be necessary to choose processing conditions that lead to intense evaporation from the surface of the particles. The large thermal gradients within the particle may also lead to partially molten particles, which lead to porosity in the coatings. Numerous onedimensional and elaborate two-dimensional models have been worked out for plasmaparticle heat transfer problem. In the case of good heat-conducting particles, the equation describing particle melting in the plasma is given by (Pauloski 1995):
telqd2h(Tg - Tp)
-
3
1/61-IdpppCp{(Tm
- 300) -k- A H f
}.
(6.3)
The left-hand side of Eq. (6.3) describes heat transferred from the flame to the particle,
Tg is the gas temperature and Tp is the particle temperature, which is assumed to be uniform in the entire volume of the particle. The right-hand side describes the heat absorbed
Thermal Plasma Processing
129
by the particle, pp and Cp being the density and specific heat of the particle respectively. A Hu is the latent heat of fusion of the material per unit volume and te the time for heating the particle from room temperature to its melting point. It is assumed that the energy transferred by radiation from the plasma is small.
6.2.4 Plasma processing systems Plasma-processing encompasses a wide range of operations including plasma-cutting and welding, plasma-melting and machining, plasma-spraying, chemical synthesis and mineral dissociation. A wide variety of plasma reactors have been developed based on DC non-transferred and transferred arc plasma torches. The various types of plasma reactors used for material processing applications have been described by Fauchais et al. (1983). The exact design of the reactor set-up depends on the specific application, but the general features remain more or less the same. The schematic of a typical plasma reactor developed in our laboratory for material synthesis is shown in Fig. 6.4. The major subsystems are plasma torch, powder-feeder, reaction chamber, torch power supply and control console. The reaction chamber consists of double walled water-cooled stainless steel segments provided with sampling and diagnostic ports. The plasma torch is mounted on one end of the reaction chamber as shown in the figure. There is also provision to collect the product and feed any reactive gas. The power sources generally employed are DC motor generator sets, transductor controlled rectifier units or silicon controlled rectifier (SCR) system. The SCR systems are currently the most extensively used power sources. A high frequency (HF) igniter to initiate the plasma is also included in the power supplies. The HF unit is connected in parallel with the DC power supply, which must be protected from it. The control console has instruments for the measurement and control of the torch inputs (viz. electrical, cooling water, gas and powder feed rates). Powder is injected into the plasma by means of the powder feeder. The turn-table-feed type powder feeder is commonly used. This ensures smooth volumetric flow of powder and has very good control over the feed rate and requires low carrier gas. 6.3. P R O C E S S I N G OF M A T E R I A L S
Plasma processing is an emerging technology that utilizes the exotic properties of the plasma medium to effect physical, chemical or metallurgical reactions to produce new materials or impart new functional properties to conventional materials. Plasma processing has been successfully applied to develop advanced ceramic coatings, synthesis of nanocrystalline materials, processing of minerals and ores, and treatment of hazardous wastes.
6.3.1 Plasma-spraying Plasma-spraying is a surface modification technique that combines particle melting, rapid solidification, and consolidation in a single process. This technique has been widely used to develop protective coatings of ceramics, alloys, and composites to enhance the surface properties of critical components operating in severe environment (Pauloski 1995,
130
P. V. Ananthapadmanabhan and N. Venkatramani 5
6
7
8
=
=
QIO
OlO
['~
Figure 6.4. Schematic of thermal plasma reactor for material synthesis: (1) cooling water; (2) control console; (3) power supply; (4) powder feeder; (5) plasma torch; (6) reaction chamber; (7) sampling duct; (8) powder collector; and (9) auxiliary instruments.
Gerdeman and Hecht 1972). In conventional plasma-spraying, the material to be coated is introduced in the plasma jet in powder form by means of a carrier gas. The powder particles, as they enter the plasma jet, are heated and melted and the molten droplets are accelerated to very high velocities, as high as 100-400 m/s. As these molten droplets strike the substrate surface, they flatten and get anchored to the surface irregularities to form an adherent coating. Plasma-spraying has several advantages over other coating processes. Because the high temperature generated by the plasma arc is far beyond the melting point of any known material, any material that melts and remains in the liquid state over a distinct range can be spray-coated. Although the plasma temperature may be in the range of 7,00010,000 K, the substrate temperature can be kept low (below 400 K), avoiding the risk of substrate distortion. The spray technique does not impose any restriction on the workpiece dimensions and large samples can be coated. The adherence of the coating to the substrate surface is predominantly mechanical in nature (Pauloski 1995). When the molten particles impact a roughened surface, they get interlocked to the surface irregularities resulting in a mechanical adherence, which is believed to be the predominant mechanism in plasma-sprayed coatings. However, depending on the nature of the substrate material, its temperature and the impinging particles, a metallurgical bond can result as in the case of nickel aluminide or molybdenum on steel substrates.
Thermal Plasma Processing
9
Figure 6.5.
,
~
,
131
-
SEM photograph of a typical sample of plasma-sprayed coating of alumina.
The coating quality is affected by a number of interdependent experimental parameters, which need to be carefully controlled. Thermal interaction of the particles with the plasma jet, the particle trajectory and the dwell time are important factors influencing coating adhesion and other properties. Particle size decides the particle trajectory and degree of melting, which controls the porosity and adhesion of the coatings. As it is practically impossible to have monosized powders, particle size in a narrow range is selected. Although the exact particle size depends on the nature of the material to be spray-coated, the recommended size for ceramic powder is 3 0 - 4 0 / z m and for metal powders 40-60 /~m.
6.3.1.1 Structure of sprayed deposits The sprayed layer has a typical lamellar structure, built up by successive layers on the material (Gerdeman and Hecht 1972). Heat transfer calculations show that freezing of the particles occurs in a few microseconds and that complete solidification occurs in perhaps 100/zs. The zone of thermal effect in the underlying material is, therefore, quite small and temperature gradients reach 105 K/cm. The wetting and flow properties of the liquid droplets are important since they influence the porosity within the coating and at the substrate interface. The surface morphology of a typical sample of plasma-sprayed alumina is shown in Fig. 6.5. The deposit shows grains consisting of melt-quenched particles with interganular and intragranular porosity, which are characteristic of plasma-sprayed coatings. The porosity of the coating is about 15-20%, which can be considerably reduced by optimizing the experimental parameters and carrying out the coating operation under reduced pressure.
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The flow and solidification of molten particles on impact with a cold surface is a difficult problem to treat theoretically because of the interaction of heat transfer and crystal growth. The very high velocity at which the various processes occur also makes direct experimental observation extremely difficult. Madejski (1976) has carried out a theoretical treatment of the impact of a molten droplet on a cold substrate, taking into account surface tension, viscosity and crystallization effects. The results show that the degree of flattening (D/d), defined as the ratio of the diameter of the flattened disc (D) to that of the original drop (d), is related to the density p, viscosity/z and impact velocity of the drop v by the following equation: D / d = 1.29(pvd/#) ~
(6.4)
Substituting appropriate values for A1203, in the velocity range of 100-400 m/s gives a value of 3-6 for the degree of flattening, which seems to be in reasonable agreement with various experimental observations (Zolotowski 1968). This suggests that at least for alumina, flattening is completed before crystallization becomes significant and the two processes can largely be treated separately. Plasma-spraying conditions are very similar to those occurring during splat-quenching of metals which has been widely studied. The cooling rate has been of particular interest and various techniques have been used to estimate the thermal conditions as a molten particle spreads and solidifies (Jones 1973). The main results of this work show that the cooling rate is largely independent of substrate material and is controlled by the interface between substrate and particle if the flattened particle has a thickness of less than a few #m. The estimated heat transfer coefficient under these conditions is of the order of 105 W m -2 K -1 . Applying this result to the case of plasma sprayed A1203 gives a cooling rate of the order of 107 K/s and a freezing time of about 10 #s. This also suggests that there is a large temperature gradient across the interface between the particle and the substrate (or particle and previously deposited layer) and the temperature of the underlying material remains relatively low. This is particularly so if the substrate is cooled during spraying by an air blast to prevent heating by plasma gases and to remove the sensible heat and heat of fusion of the sprayed particle. 6.3.1.2 Plasma spray-deposited materials Applications for plasma-sprayed materials range from microelectronics to large engineering structures. There are excellent reviews and publications (Pauloski 1995, Gerdeman and Hecht 1972, Zaat 1983, Fauchais and Vardelle 1994, Notomi and Hisatome 1996, Boulos 1992, Matejka and Benko 1989) exclusively devoted to the coating process and applications, which can be referred to for further details. Plasma-sprayed coatings of metals, alloys and ceramic materials have been used for various applications. Plasma-spraying has been applied to produce wear-resistant coatings of ceramics and metals on various components. Successful wear-resistant coatings have been developed using cobalt- and nickel-based alloys containing carbides of chromium, tungsten and titanium as abrasion-resistant particles (Rangaswamy and Herman 1986, Mohanty and Smith 1995, Knight et al. 1995). Self-lubrication can be provided by incorporating low-friction materials or dry lubricants such as graphite, molybdenum disul-
T h e r m a l P l a s m a Processing
133
phide, or organic lubricants. Plasma-spraying has been effectively used to protect metallic and non-metallic components from high temperature and oxidation. Plasma-sprayed thermal barrier coatings consisting of multi-layered and graded ceramic and metal coatings are being extensively used for protecting the surfaces of metal components, gas turbines and other critical engine parts exposed to corrosive environment. Another important application of plasma-spraying is reclamation, where it is used to repair worn out or badly machined components. Worn out surfaces of brake drums, grinders, spindles, etc. can be reclaimed by plasma coating. Oxide coatings are used for many high-tech as well as conventional applications. Plasma-sprayed coatings of stabilized zirconia, alumina, and alumina-titania have been extensively studied (Bratton and Lau 1981, Stecura 1977, Ananthapadmanabhan et al. 1991, Ramachandran et al. 1997). Aluminium oxide (often mixed with titania) and chromium oxide are used on many textile machinery components such as thread guides (Matejka and Benko 1989). They are also widely used on pumps and mechanical seals to provide resistance to both abrasion and corrosion. Plasma-sprayed zirconia coatings, as thermal barriers, are effective in improving the efficiency of gas turbine engines. The preferred material is 6-8 wt.% Y203-ZrO2 or 24 wt.% MgO-ZrO2 deposited over Ni-CrA1-Y bond coat applied by vacuum plasma-spraying (Meetham 1988). The latter provides oxidation resistance to the superalloy substrate and is also strain tolerant, which helps to accommodate stresses arising out of thermal expansion mismatch between the base metal and the ceramic coating. Thermal shock resistance of zirconia coatings has been greatly enhanced by using composite coatings of alumina and calcia stabilized zirconia. The improved thermal shock behavior is attributed to destabilization of calcia stabilized zirconia leading to the generation of microcracks which act as thermal stress relief centers (Ananthapadmanabhan et al. 1991). When ceramic coatings are applied on metal substrates, the large mismatch in the thermal expansion coefficients of the metal base and the coating material can result in interface stresses. This can lead to poor adhesion and even coating failure. In order to overcome this, an intermediate coating, referred to as bond coat, is applied on the metal substrate before applying the ceramic coating. The primary function of the bond coat is to reduce the interface stress and improve coating adhesion (Matejka and Benko 1989). Nickel aluminide and its derivatives, Ni-Cr-A1-Y, are frequently used as the bond coat material. The reaction of nickel and aluminum to form the intermetallic is highly exothermic, resulting in high adhesion strength. Nickel aluminide coatings are usually prepared by using special composite powder of nickel and aluminum (Matejka and Benko 1989). This composite powder forms the compound during the spray process. A one-step process to form plasma-sprayed Ni3A1 coatings has been developed (Ananthapadmanabhan et al. 1995) by using a blend of nickel and aluminum powders as the feedstock material. The process parameters have been optimized to effect compound formation as the powder particles deposit on the substrate surface. Coating adhesion was measured by tensile adhesion strength as described in ASTM C 633 (Lin and Berndt 1994). The adhesion strength of the coatings was found to be much greater than 60 MPa.
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The cooling rate for plasma-sprayed particles approaches 106 K]s and this produces a fine-grained microstructure within the splat (Scott and Kingswell 1991). This rapid solidification can produce amorphous deposits from some ferrous alloys (Herman and Bhat 1980), while metastable or non-stoichiometric phases can be retained in some ceramics. Phase composition and stability of yttria stabilized zirconia coatings have been extensively studied. The phase structure of yttria stabilized zirconia coatings have been shown to consist of non-equilibrium phases. The work of Miller et al. (1981) and Bratton and Lau (1981) on plasma-sprayed zirconia thermal barrier coatings showed that the coatings consisted of non-equilibrium phases. Irrespective of the starting material, the as-sprayed material was found to contain a non-transformable tetragonal (high yttria) phase. It was observed that partially stabilized zirconia (ZRO2-8% Y203) as well as fully stabilized zirconia (ZRO2-20% Y203) contain primarily the non-equilibrium tetragonal phase, which transformed to the equilibrium cubic and tetragonal (low yttria) during thermal annealing. Plasma-sprayed coatings of alumina and alumina-titania have been shown to consist of metastable alumina (Heintze and Uematsu 1992, Ramachandran et al. 1997). The work of Heintze and Uematsu (1992) has shown that plasma-sprayed coatings of alumina predominantly consist of y-alumina. They also found that u-alumina can be formed by careful optimization of the process parameters.
6.3.1.3 Low-pressure plasma-spraying Atmospheric plasma-spraying (APS) is seldom satisfactory when used to prepare coatings of metals, alloys, and non-oxide ceramics. When coatings of refractory carbides, nitrides and borides are prepared by APS, the high velocity plasma jet acts as an aspirator and sucks in air from the surrounding atmosphere resulting in oxide inclusions, which can seriously affect the coating performance. Studies on the phase composition of TiB2 coatings prepared by APS showed appreciable amount of boric oxide, formed by partial in-flight oxidation of TiB2 particles (Ananthapadmanabhan et al. 1993). The low melting point of B203 (about 800 K) affects the high temperature performance of TiB2 coatings. This problem can be overcome by low-pressure plasma-spraying (LPPS) or vacuum plasma-spraying (VPS), where the coating operation is performed in a vacuum chamber. Alternately, the chamber can be filled with inert gas to minimize coating contamination. The quality of the coatings produced by LPPS is superior to that of those prepared by APS, in terms of coating properties as well as chemical purity. LPPS reduces porosity and improves coating adhesion. Reactive metals like titanium, zirconium and tungsten can be spray-coated with minimum oxide contamination. LPPS has been extensively used to prepare coatings of refractory carbides, nitrides and other non-oxide materials for various applications. Recent developments include the successful application of LPPS to fabricate components for solid oxide fuel cell (SOFC) components (Notomi and Hisatome 1996). Dense electrolyte coatings of yttria stabilized zirconia have been prepared by low pressure plasma-spraying. A 1 kW module has been developed and successfully operated for 3,000 h.
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6.3.1.4 Reactive plasma-spray f o r m i n g Conventional plasma-spraying consists in feeding metal, alloy, or ceramic powders in the plasma jet, where the particles melt and the molten droplets are deposited on to a substrate surface, forming an adherent coating. It is essentially a physical process involving melting and resolidification of the powder. In reactive plasma-spraying, the high temperature plasma environment is utilized to realize inflight chemical reactions by using suitable reactive gases along with the feed powder (Mutasim et al. 1992). The metal or ceramic powders are introduced in the plasma jet along with a suitable reactive gas. The reactant powder particles melt and the molten droplets react with the reactive gas, resulting in the in situ formation of the desired compound, which is deposited on the substrate. This technique has been used to develop refractory carbide coatings. Reactive plasma-spraying of Ti using methane, propylene, or acetylene as the reactive gas has been carried out (Mutasim et al. 1992, Ananthapadmanabhan et al. 1993) to produce Ti-TiC composite coatings. Dissociation of the hydrocarbon to form finely dispersed aerosol of carbon is the first step in the formation of the carbide. These fine particles of carbon subsequently react with the molten droplets of the metal to form the metal carbide, which is deposited on the substrate. By proper selection of the process parameters, composite coatings containing 60 vol.% of TiC dispersions have been obtained. It was also observed that higher substrate temperatures increased the carbide content. A recent development in this field is a one-step process to form Fe-TiC cermet coatings by reactive plasma-spraying of ilmenite (Ananthapadmanabhan and Taylor 1998). The principle involves the reduction of ilmenite by methane in a thermal plasma environment according to the following chemical reaction:
FeTiO3 + 4CH4 = Fe + TiC + 3CO + 8H2.
(6.5)
Ilmenite powder is introduced into the plasma jet using a mixture of methane and argon as the carrier gas and the substrate is kept at about 1200 K. Ilmenite is reduced to iron and titanium dioxide, which is further reduced to its lower oxides and finally to TiC. Fig. 6.6 shows the surface morphology of a typical Fe-TiC coating, revealing excellent inter-particle cohesion and very little porosity. Reactive plasma-spraying can be extended to develop other refractory carbide and nitride coatings by suitable choice of precursor materials and reactive gases. Reactive plasma-spraying is also used for depositing hydroxy apatite (HA) coatings for dental implants. Conventional atmospheric plasma-spraying of hydroxy apatite, Cal0(PO4)6(OH)2 results in the formation of biodegradable phases such as, fi-tricalcium phospate and tetracalcium phosphate by the thermal dissociation of HA, rendering it unsuitable for biomedical applications (Hench 1991). The decomposition of HA to biodegradable compounds has been eliminated by a reactive plasma-spray technique developed by Patil et al. (1996).
6.3.1.5 Plasma spheroidization Powders of metals, alloys and ceramics can be melted in a plasma jet. Melting of a particle results in the formation of a spherical drop under the action of the surface tension forces and this shape is usually retained after solidification, thus providing the name to
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Figure 6.6. SEM photograph of Fe-TiC coating formedby reactive plasma-spraying.
the process. This treatment may be used simply to give a spherical particle shape for particular application, such as Fe304 for photocopying, plasma spray-quality powders for thermal spray applications and UO2 for dispensed nuclear fuels. Spheroidization is particularly useful to prepare spray-quality powders of special materials for thermal spray applications. Spheroidization of a powder blend of Ni-15% A1 has been carried out in a thermal plasma reactor (Ananthapadmanabhan et al. 1997). The as-collected powder showed poor crystallinity as indicated by the broad X-ray diffraction pattern (Ananthapadmanabhan et al. 1995). Annealing at 800 K in an atmosphere of flowing argon resulted in the formation of well-crystalline Ni3A1. The process parameters have been optimized to yield a precursor powder of Ni-A1, which forms Ni3A1 on the substrate surface during plasma-spraying. A typical photomicrograph of plasma spheroidized Ni-A1 powder is shown in Fig. 6.7. Besides possessing spherical morphology, the powder has excellent flow characteristics, making it ideally suited for thermal-spray applications. Spheroidized particles may also be prepared by atomization of rods or wires fed into a plasma. The diameter of the particles produced depends on the diameter of the torch nozzle, the gas flow rate, plasma density and temperature and surface tension of the material in the liquid state (Rykalin 1976). Alumina particles of 40 # m diameter have been produced by feeding alumina rod in a hydrogen plasma. The heat transfer can be improved if the arc is transferred to the wire but this is not possible with non-conducting materials. An example of large scale industrial application is the spheroidization of magnetite for photocopying applications, where magnetite particles, 125/zm diameter, are produced in 600 kW AC air plasma heater with a power consumption of 2 kWh&g (Fey et al. 1975).
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Figure 6.7. SEM photograph of plasma spheroidized Ni-A1 powder.
6.3.2 Plasma reactors and furnaces
Plasma technology can be gainfully utilized to treat minerals and ores to yield value-added materials. Decomposition processes which are already proved both thermodynamically and technically feasible and favorable include dissociation of carbonates, oxides, sulfides, halides and complex polymetallic ores (Dembovski 1985). It has also been used for metallurgical applications and in the synthesis of advanced ceramics. Table 6.4 shows the high power plasma systems used for diverse applications.
6.3.2.1 Plasma decomposition Decomposition of zircon sand in thermal plasma reactors to produce zirconia is well established. The principle involved is the thermal dissociation of zircon into its component oxides, followed by quenching the reaction products to yield a mixture of ZrO2 and SIO2. Zircon is thermodynamically stable up to 1950 K, above which it spontaneously dissociates into zirconium dioxide and silicon dioxide. The reverse reaction is prevented by freezing the high temperature equilibrium by rapid quenching of the products. Zircon sand is fed into the plasma by means of a powder feeder. As the zircon particles pass through the plasma, they are heated to temperatures above 2,000 K, when dissociation takes place and the product is rapidly quenched. Zirconia is separated from the oxide mixture by leaching away silica with 30-50% solution of sodium hydroxide. Extensive studies on the influence of operating parameters on the dissociation reaction have been carried out by Ananthapadmanabhan et al. (1994a). Results of the study showed that the extent of dissociation of zircon was affected by the power input to the
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P. V. Ananthapadmanabhan and N. Venkatramani
Table 6.4 High-powerplasma torches for material processing applications Torch make Huls Linde-Retech Tioxide Aero-Spatiale SKF Ionarc Metco Daido Tetronics Voest-Alpine Chemische Werke Huls AG Voest-Alpine
Power ( M W ) 8.5 0.7 5.0 5.0 7.0 0.35 0.1 1.0 3.0 7.5 8.0 4.0
Applications Chemical synthesis Plasma melting Titanium dioxide Ferro-manganese Recovery from dust Zirconia production Plasma-spraying Scrap melting Tundish heating Steel remelting Synthetic gas for direct reduction of iron ore Ferromanganese
plasma, plasma gas flow rate, carrier gas flow rate, feed rate of zircon, etc. Ionarc Tafa has developed a 300 kW multi-torch furnace for processing zircon sands at 300 kg/h feed rate. The process has been shown to be economically superior to the conventional process. The product purity is superior with an electricity consumption of 1.32 kWh/kg of the product compared to 9.9 kWh/kg for the conventional process (Dembovski 1985). A plasma reactor has also been used to convert ilmenite to titanium dioxide. The process involves reductive dissociation of ilmenite to TiO2 and iron in a thermal plasma medium using controlled amount of methane as the reductant. The process has been extended to prepare Fe-TiC powder (Taylor et al. 1995).
6.3.2.2 Plasma metallurgy The use of RF and DC plasmas for metallurgical applications is getting established in many countries throughout the world. The introduction of plasma furnaces has revolutionized the metallurgical industry by improving the operational efficiency and quality of the products. The conventional iron and steel industry is highly capital-intensive. The blast furnace operation utilizes dwindling resources of expensive high-grade coking coal, which is used for heating and reducing operations. Consumption of this expensive coke can be considerably reduced by augmenting the conventional blast furnace operation with the plasma process (Dembovski 1985). The process consists in partially replacing coking coal by superhot reducing gas synthesized by plasma-heating of a mixture of natural gas and air or carbon dioxide. Superhot reducing gas can reduce the use of coke by as much as 75%, resulting in substantial cost savings. This process is referred to as the direct reduction process and the direct-reduced iron (DRI) is used for producing low-carbon steels. Mintek, South Africa has developed pilot-plant-scale furnaces for producing low-carbon steel by smelting DRI and iron ore (McRae et al. 1985).
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The application of thermal plasma technology to steel-plant dust is a recent development. Dust and metal-bearing wastes from steel plants consist of oxides of various metals and their disposal in landfill sites poses environmental hazards, since toxic metals may be leached into drinking water supplies. Thermal plasma processing of steel-plant dust not only offers an effective way to destroy the hazardous waste, but also to recover valuable metals such as zinc, nickel and chromium. SKF of Sweden has developed the Plasmared processes for recovering metals from metal-bearing wastes in a reducing atmosphere generated by the decomposition of fossil fuels in a plasma stream. The Enviroplas process is a DC plasma arc process, developed by Mintek, South Africa, which can treat arc-furnace dust, alloy-steel dust and other metal-bearing wastes. Production of ferrochrome, ferromanganese and nitrogen alloying of steel are other areas where plasma technology has been extensively used. DC transferred arc furnaces have replaced the conventional submerged arc furnaces for producing ferrochromium from chromite (FeCrO3) ore. The DC arc process ensures higher chromium recovery with reduced levels of phosphorous and sulfur (Jones et al. 1997). Nitrogen alloying of steel by plasma treatment is illustrative of the unique advantages of the process over conventional metallurgical alloying processes. Interstitially dissolved nitrogen improves the mechanical properties and corrosion behavior of steel significantly. The strength and corrosion resistance of some steel grades are commonly enhanced by addition of nitrogen-enriched ferrochrome, ferromanganese or calcium nitride. In conventional induction furnaces with nitrogen partial pressure of 105 Pa, the amount of nitrogen dissolved in steel at 2,500 K is hardly 0.06%, which can be increased to about 0.2% by weight under very high partial pressures of nitrogen. In plasma furnaces, however, a substantial fraction of nitrogen is present in atomic, excited and ionic state. Under identical conditions of temperature and pressure, plasma processing can increase the dissolved nitrogen content in iron or steel by as much as an order of magnitude over the levels attained in induction furnaces. The tensile strength, yield point and creep strength of 18Cr-8Ni steel are found to increase with the amount of dissolved nitrogen. A nitrogen content of 0.8% raises the yield strength from about 200-700 MPa (Dembovski 1985). Substantial improvement in the long-term creep strength and corrosion resistance in various corrosive media are also observed. Reduction of metal halides and oxides to metal in hydrogen plasma offers many technological advantages. The atomic hydrogen and other active species present in the plasma medium considerably enhance the kinetics of the reduction reaction. The high residual hydrogen contents of metals gained by this technique can be brought down to acceptable levels by subsequent vacuum annealing. Zirconium tetrachloride vapor has been reduced with hydrogen to produce metallic zirconium in a hybrid DC transferred arc plasma reactor (Jones et al. 1997). This plasma reduction technique has been successfully applied to reduction of WO3 and MOO3. Arc plasma devices operating in the range of 150-250 kW have been used to produce tungsten and molybdenum. Similar oxide reduction processes have been extended to produce metals of Ni, Cr and other elements using an argon plasma medium containing additions of hydrogen and hydrocarbons.
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6.3.2.3 Processing o f ceramics High-temperature vapor phase chemistry in the thermal plasma medium offers an attractive route for powder synthesis. The process offers flexibility in the choice of the reactants, which can be solids, liquids or gases. The starting material is introduced in the form of a gas, liquid or solid. The reaction proceeds in several steps: vaporization, vapor phase reaction, nucleation and condensation. Carbides, nitrides, borides and oxides of refractory metals have been successfully produced in RF and DC torch plasma-based reactors. 6.3.2.3.1 Oxide powders. Plasma processing has been widely used to produce ultrafine powders of alumina, magnesia, titania and other oxide powders. A plasma-based technology has been developed for large-scale manufacture of pigment-grade titanium dioxide by the Tioxide Group in UK under the name ICON process. The process essentially makes use of a DC plasma torch to oxidize TIC14. The principal advantages of the plasma process over the conventional process are simplification of the process steps, narrow particle size range and precise control over product quality. Ultrafine silica is used as a filler material in composites and as a thixotropic additive material in the paint industry. The active silica product is prepared by vaporizing SiO2 in an oxidizing atmosphere followed by rapid quenching. The resulting powder is highly active and amorphous. Ultrafine aluminum oxide, iron oxide and zirconium oxide have been synthesized in an inductively coupled plasma (ICP) reactor (Kagawa et al. 1983). Nitrate solutions of the respective metals were introduced into the plasma jet, where decomposition of the fine droplets of the nitrate into the oxide occurred. The fast quench rate favored the formation of metastable ),-phase of A1203 from aluminum nitrate. The dissociation of zirconium nitrate in the ICP reactor resulted in the formation of ZrO2, which was predominantly in the metastable tetragonal phase. 6.3.2.3.2 Non-oxide ceramics. The carbides of silicon, titanium, boron, and tungsten are characterized by high hardness and excellent wear and corrosion resistant properties. These materials find application in the cutting tool industry and also as abrasives. The high energy density of the transferred arc has been used to produce metallic vapors that on condensation produce submicron aerosols. These vapors are also reacted in a suitable environment to produce a variety of products. Submicron size silicon carbide powders have also been produced with SIC14 and methane in a DC plasma torch (Allaire et al. 1991). Ultrafine powders of silicon nitride and aluminum nitride have also been produced using arc plasmas using ammonia as the nitriding gas (Ageorges et al. 1994). Processing of ultrafine powders of refractory carbides and nitrides is one of the widely studied areas of plasma processing. Most of the research efforts have been directed to the study of gas-phase reactions in plasma medium involving metal halides and suitable reactive gas to give ultrafine powders of refractory carbides and nitrides (Yoshida et al. 1979). Titanium carbide and tungsten carbide have been synthesized in DC and RF plasma reactors from titanium tetrachloride in a methane atmosphere. Ti and W metal powders have also been used to synthesize TiC and WC powders (Ananthapadmanabhan et al. 1994b). Plasma processing in RF plasma torches has also been used to prepare titanium
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141
nitride. The process involves the use of titanium halide or titanium metal powder with ammonia or nitrogen as the reactive gas. The liquid injection plasma synthesis (LIPS), where a liquid precursor is injected into the plasma jet, is widely used to process nanosized ceramic powders (Kong and Lau 1990). The liquid injection can be either parallel to the flow or counter current to the flow. Ultrafine oxide and non-oxide ceramic powders and mixed oxide compounds have been synthesized using the LIPS and counter current plasma synthesis (CCPS). Refractory carbides, nitrides, high temperature super-conducting oxide and mixed oxide powders have been produced in RF and DC based plasma reactor systems. 6.3.2.4 Treatment of hazardous wastes Conventional incinerators are, often, unsatisfactory for effective treatment and disposal of hazardous wastes, which are often mixtures of metals, organic and inorganic compounds. Production of toxic gases such as dioxin, insufficient thermal destruction of inorganic constituents, production of polluted flue gases, etc., are some of the drawbacks of conventional waste management systems. Thermal plasma technology has been shown to be potentially useful in the management of industrial waste (Boulos 1996), hospital waste, etc. Multi-torch plasma furnaces have been developed for destruction of highly toxic waste materials such as polychlorinated biphenyls (PCBs), pesticides, herbicides and organic solvents. These systems are also suitable for nuclear waste treatment (Ageorges et al. 1994).
6.3.3 Processing of metastable phases The unique feature of plasma processing is its ability to overcome thermodynamic and kinetic barriers and gain access to non-equilibrium regimes. The high temperature and the high quench rate (105-107 K/s) help in processing metastable and amorphous phases. As the reactants enter the plasma jet, they are heated to their melting temperatures and very often, an appreciable fraction of the particles vaporize and react with each other. The atomic and molecular clusters of the product are carried to the cooler regions of the reactor, where they are quenched. The entire process is completed in a few milliseconds. This time frame is too small for atomic rearrangement to form crystalline structures. It has been observed that thermal plasma processing of silica and other silicate minerals, such as zircon invariably produces highly reactive silica in an amorphous state as confirmed by X-ray diffraction of the product. Plasma process is capable of producing metallic glasses, amorphous silicon, diamond and other metastable phases of ceramic materials.
6.3.3.1 Plasma deposition of diamonds Conventionally, diamond is synthesized by recrystallization of metal-solvated carbon under high temperature (2,000~ and extreme pressures (50-100 kbar). Vapor-phase synthesis of diamond by the hot-filament chemical vapor deposition (CVD) technique has been known for many years. However, the technique is characterized by extremely low deposition rates. The use of a thermal plasma jet for diamond deposition is a recent development. Thermal dissociation of methane and other hydrocarbons in DC and RF plasma jets have been successfully applied to develop diamond coatings (Yoshida et al.
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P. V. A n a n t h a p a d m a n a b h a n a n d N. Venkatramani
Figure 6.8 Surfacemorphologyof plasma deposited diamond
1979, Meyer et al. 1989, Metaxas 1996). In contrast to typical growth rates of 0.05/xm/h in hot filament CVD methods, the plasma process is characterized by high growth rates of the order of 900 #m/h (Ohtake et al. 1991). Ramachandran et al. (1996) have injected a mixture of methane and hydrogen gas into an argon plasma jet issuing out of a DC arc plasma torch operated in the non-transferred arc mode. Figure 6.8 shows the surface morphology of diamond deposited on silicon substrate. The cauliflower-like structure is characteristic of high growth rate. The Raman spectrum, indicating the characteristic Raman Shift of the diamond deposit, is shown in Fig. 6.9. The precursor gases employed consisted of a mixture of methane and hydrogen in the ratio 1:99. The experimental conditions strongly influence the diamond growth rate, both qualitatively and quantitatively (Baldwin et al. 1994). In particular, the substrate temperature and hydrogen/methane ratio are significant. Although the exact mechanism of diamond formation is still not understood, high temperature gas phase reactions under plasma environments involving methyl radicals is proposed to be the main operating mechanism.
6.3.3.2 Thermal plasma synthesis of ultrafine alumina powder A classic example illustrating the non-equilibrium nature of reactions occurring in the plasma medium is exemplified by plasma oxidation of aluminum powder (Ananthapadmanbhan et al. 1996). The experimental set-up used is shown in Fig. 6.4. Aluminum metal powder, stored in a powder feeder, was fed into the plasma using argon as carrier gas. Compressed air was injected into the reactor. The powder particles melt and vaporize and the molten droplets and vapors of aluminum react with oxygen, resulting in the
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143
800
4330
cr~ t
600
400 ,,
A
15IXl
Wovenumbers ( cn~ I )
|
1000
Figure 6.9 Raman spectrum of diamond deposit.
Table 6.5 Typicalexperimental parameters Power input (kW) Arc voltage (V) Plasma gas Ar (LPM) Secondary gas N2 (LPM) Carrier gas Ar (LPM) Reactive gas air (LPM) Powder feed rate (g/min)
16 40 20 02 10 40 05
formation of aluminum oxide that is deposited on the reactor walls. Typical operating parameters are given in Table 6.5. X-ray diffraction results indicate that the powder consists of the metastable ),aluminum oxide. Formation of the ),-phase is consistent with the highly non-equilibrium nature of the process. Vapors and molten droplets of aluminum react with oxygen to form aluminum oxide, which crystallizes in the metastable ),-phase by virtue of the rapid quench rate associated with the process. Figure 6.10 is a scanning electron micrograph (SEM) of a typical sample of alumina powder collected from the main section of the reactor. Individual particles are not discernible, but agglomerates consisting of nanosized particles can be seen. A transmission electron micrograph (TEM) of a typical sample is shown in Fig. 6.11. The clusters are resolved into individual particles with a size ranging
144
P. V. Ananthapadmanabhan and N. Venkatramani
Figure 6.10. SEM photographof plasma synthesized alumina
from a few nanometers to about 30 nm. The spherical morphology of the particles is also evident from the figure.
6.3.3.2.1 Formation of },-alumina. The free energy minimization plot for the system A1O-N has been developed using the CSIRO thermopackage (Turnbull and Wadsley 1988). Figure 6.12 shows the thermodynamic equilibrium diagram for the system aluminumoxygen-nitrogen in the molar ratio 1:2:8 (close to the experimental conditions) under 1 atmosphere pressure. The Y-axis shows the concentration of the species in number of mols and the X-axis shows the temperature. Argon as the plasma gas is also included in the thermodynamic calculations. However, only the oxides of aluminum, atomic species of oxygen and nitrogen and aluminum vapor are shown in the figure. It can be seen from Fig. 6.12 that at temperatures greater than 5,000 K, which are commonly encountered in plasma reactors, the stable species present are aluminum vapour, atomic oxygen and atomic nitrogen. In the temperature region 3,500-4,500 K, A10(g) becomes increasingly stable. The other oxide species of aluminum such as A120, A102 and A1202 are also formed in this temperature zone, although at lower concentration levels, possibly by the following reactions: Al(g) + O(g) = A10(g)
(6.6)
145
Thermal Plasma Processing w
~
" ,'~t~" ",
. . ~.. .
.:,:
*
'
.:
~
.
.
..
'*~" , ~
~--
.
.
100
nm
Figure 6.11. Transmission electron micrograph of plasma synthesized alumina.
2Al(g) + O(g) = A120(g)
(6.7)
2Al(g) + 20(g) = A1202(g)
(6.8)
Al(g) + 20(g) = A102(g)
(6.9)
2A10(g) = A120(g) + O(g).
(6.10)
As the temperature falls below 4,500 K, oxygen molecules and nitrogen molecules (not shown in the figure due to their high concentration) start forming by the recombination of the respective atomic species. At temperatures below 3,500 K and above 2,500 K, liquid aluminum oxide is formed. Solid ot-A1203, which is the only thermodynamically stable form of alumina, results as the temperature drops to below 2,300 K, which is the
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P. V. Ananthapadmanabhan and N. Venkatramani
1 0.9 0.8 0.7 0.6 o.s
~s)
N(g)
0.4 0.3 0.2 0.1 0 0 0 0 '~--
0 0 it) ~--
0 0 0 tN
0 0 rid r
0 0 0 03
0 0 ur) ~
0 0 0 ~I*
0 0 it') ~
0 0 0 IZ)
0 0 it') it')
0 0 0 t,,ID
T ~ ~ Figure 6.12. Free energy minimization plot for the system AI-O-N (1:2:8).
melting point of alumina. Below 2,300 K, the stable species are solid ot-A1203, nitrogen and oxygen molecules (not shown due to their high concentration). Thermodynamic considerations would lead us to conclude that ct-Al203 is formed. However, the conditions existing in the plasma reactor, augmented by thermochemical effects associated with the oxidation of aluminum, favor the formation of non-equilibrium phases. The first step leading to the formation of alumina is the melting and vaporization of aluminum powder injected into the plasma. Molten droplets and aluminum vapor react with oxygen downstream of the reactor to produce aluminum oxide according to the following equation: 2A1 + 3 0 2 -- A1203, AH ~ -- - 1675.7 kJ/mol.
(6.11)
Oxidation of aluminum to alumina is highly exothermic, the standard enthalpy of formation of aluminum oxide from the elements being -1675.7 4- 1.2 kJ/mol (Chase et al. 1985). The enthalpy of formation increases with the reaction temperature and is -2243.73 kJ/mol at 3,000 K. The heat released during oxidation coupled with the high temperature in the plasma medium vaporizes aluminum oxide, leading to the formation of suboxides of aluminum and if the temperature is higher than 4,500 K aluminum vapor and oxygen atoms are formed. As aluminum vapor and oxygen atoms move down to regions where the temperature is between 3,500-4,500 K, they react to form A10 gas and other oxides of aluminum. This is followed by a rapid quench of the vapors as they move to the cooler zones of the reactor tube. The quench rate in plasma processing has been estimated to be
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of the order of 106-108 K/s and the extent of supercooling achievable is much greater than 0.2Tm, favoring homogeneous nucleation resulting in nanosized particles of ),-alumina. In principle, the selection of a metastable phase during solidification is associated with kinetic factors and requires suppression of the nucleation of the stable phase. (See also Chapter 2 on "Thermodynamics and Kinetics of Metastable Phase Formation" by K. N. Ishihara in this book). The nucleation rate of a solid condensing from liquid or gaseous phase is decided by its critical nucleation free energy, A G*, which is the energy barrier for nucleation. The relative nucleation rates of a and y phases of alumina depend on the ratio of their respective critical nucleation free energies, A G * / A G•*. McPherson (1973) has estimated the critical free energies of nucleation for a- and v-alumina as a function of temperature. The estimated ratio of the relative critical free energy for nucleation of a- and y-alumina indicates that v-alumina would nucleate in preference to a-alumina at temperatures lower than 1,740~ The estimated solidification temperature for alumina particles falls in the range of 1,500-1,600~ under plasma conditions, suggesting the formation of v-alumina in preference to a-alumina. Formation of v-alumina in preference to the thermodynamically stable a-phase is typical of rapid quench techniques such as plasma-spraying (Heintze and Uematsu 1992), arc discharge (Hirayama 1987) and electro-hydrodynamic atomization (Jayaram and Levi 1989). Spherical particles of alumina prepared by electro-hydrodynamic atomization have been shown to consist of metastable phases like y- and 6-alumina. Aluminum oxide prepared by vapor condensation from arc, flame oxidation of aluminum salts also consists of metastable phases rather than the thermodynamically stable form. However, the final phase obtained depends on the thermal excursion the particles experience during recalescence and subsequent cooling of the solid powder. This depends on the particle size and the cooling rate. The transformation of y-A1203 to a-A1203 for plasma-sprayed alumina droplets has been analyzed by McPherson (1981), who estimated the achievable supercooling and the thermal history of droplets of different sizes. The analysis shows that particles less than 1 # m in diamater consist entirely of y-A1203, whereas particles >50 # m are completely transformed to the a-phase. Levi et al. (1988) have developed a more rigorous model to examine microstructure evolution continuously through the liquid cooling stage, recalescence and subsequent cooling of the solid powder. Their analysis shows that decreasing the droplet size results in higher supercoolings and more efficient cooling leading to the formation of the metastable phase. The particle size range normally observed in plasma processed alumina is of the order of a few tens of nanometers suggesting that high supercoolings favor the formation of the v-phase. The above analysis helps to illustrate the physicochemical principles involved in the formation of non-equilibrium phases under plasma environment. The generalities of the model can be extended to any other specific case. 6.4. S U M M A R Y AND C O N C L U S I O N S
An overview of the recent developments in the field of thermal plasma processing has been presented. Thermal plasma processes are characterized by high temperature, high enthalpy and power density, 10-100 times higher than those encountered in conventional
148
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a n d N. V e n k a t r a m a n i
furnaces. Different plasma devices ranging from a few hundreds of k W to a few M W p o w e r levels have been designed and developed, over the past decade, to meet a variety of requirements ranging from p o w d e r synthesis to primary and secondary metallurgical operations and treatment of hazardous wastes. The use of plasma-sprayed ceramic and alloy coatings for industrial applications is a well established technique. But, there have been newer developments such as the reactive plasma-spray technique, which combine plasma processes and chemical reactions in one single operation. A further development in this area is the direct conversion method using cheap raw materials. The gas-to-solid conversion in plasma process to fabricate diamond coatings is a variant of this method. Contrary to conventional processing techniques, which are limited by thermodynamic and kinetic barriers, plasma processing, by virtue of the high temperature, high power density and rapid quench, offers unlimited scope for preparation of metastable phases and amorphous structures. The presence of reactive species provides new reaction paths to synthesize a variety of novel materials.
ACKNOWLEDGMENTS
The authors thank Mr. U. K. Chatterjee, Head, Laser & Plasma Technology Division for his constant support and encouragement. The cooperation and help received from Mr. K. P. Sreekumar and Dr. D. S. Patil are also gratefully acknowledged.
REFERENCES
Ageorges, H., Megy, S., Chang, K., Baronnet, J. M., Williams, J. K. and Chapman, C. (1994) Plasma Chem. Plasma Process., 13, 613. Allaire, E, Parent, L. and Dalleuie, S. (1991) J. Mater. Sci., 26, 4160. Ananthapadmanabhan, P. V. and Taylor, P. R. (1998)J. Alloys and Compounds (in press). Ananthapadmanabhan, P. V., Sreekumar, K. P., Muraleedharan, K. V. and Venkatramani, N. (1991) Surf. Coat. Technol., 49, 62. Ananthapadmanabhan, P. V., Sreekumar, K. P., Ravindran, P. V. and Venkatramani, N. (1993) Thin Solid Films, 224, 148. Ananthapadmanabhan, P. V., Sreekumar, K. P., Iyer, K. V. and Venkatramani, N. (1994a) Mater Chem. Phys., 38, 15.
Ananthapadmanabhan, P. V., Sreekumar, K. P. and Venkatramani, N. (1994b) Plasma Chemical Processing of Refractory Carbides, Plasma Science and Technology, ed. Das, I. M. L. (Allied Publ., New Delhi, India), p. 627. Ananthapadmanabhan, P. V., Sreekumar, K. P., Ravindran, P. V. and Venkatramani, N. (1995) in Thermal Spraying--Current Status and Future Trends, Vol. 2, ed. Ohmori, A. (High Temperature Soc. of Japan, Tokyo), p. 1127. Ananthapadmanabhan, P. V., Sreekumar, K. P., Venkatramani, N. and Taylor, P. R. (1996) J. Alloys and Compounds, 244, 70. Ananthapadmanabhan, P. V., Sreekumar, K. P., Venkatramani, N. and Taylor, P. R. (1997) in EPD Congress 1997, ed. Mishra, B. (TMS, Warrandale, PA), p. 209. Baldwin, Jr., S. K., Owano, T. G. and Kruger, C. H. (1994) Plasma Chem. Plasma Process., 14, 383. Bonet, C. (1976) Chem. Eng. Prog., 72, 63. Boulos, M. (1992) J. Thermal Spray Technol., 1, 33. Boulos, M. I. (1996) Pure Appl. Chem., 68, 1007. Bratton, R. J. and Lau, S. K. (1981) inAdvances in Ceramics, Vol. III, Science and Technology ofZirconia, eds. Heuer, A. H. and Hobbs, L. W. (American Ceramic Society, Columbus, OH), p. 226.
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Chase, Jr., M. W., Davies, C. A., Downey, Jr., J. R., Frurip, D. J., McDonald, R. A. and Syverud, A. N. (eds.), (1985) JANAF Thermochemical Tables, third edition (The American Chemical Society and The American Institute of Physics for the NIST, New York). Dembovski, V. (1985) Plasma Metallurgy: The Principles, Materials Science Monographs, No. 23, (Elsevier, Amsterdam). Eliezer, Y. and Eliezer, S. (1989) The Fourth State of Matter: An Introduction to the Physics of Plasma (Adam Hilger, Bristol, UK). Fauchais, E and Vardelle, M. (1994) Pure Appl. Chem., 66, 1247. Fauchais, P., Bourdin, E., Coudert, J. E and McPherson, R. (1983) in Topics in Current Chemistry, Vol. 107, eds. Veprek, S. and Venugopalan, M. (Springer-Verlag, Berlin/Heidelberg), p. 59. Fey, M. G., Wolf, C. B. and Harvey, J. E (1975) in Proceedings of the International Round Table: Transport Phenomena in Thermal Plasmas (Odeillo, France). Fineman, J. (1987) Plasma Technology in Metallurgical Industry (Iron and Steel Society, Warrendale, PA). Gerdeman, D. A. and Hecht, N. L. (1972) Arc Plasma Technology in Materials Science (Springer-Verlag, Wien). Heberlein, J. V. R. (1992) Pure Appl. Chem., 64, 629. Heintze, G. N. and Uematsu, S. (1992) Surf. Coat. Technol., 50, 213. Hench, L. L. (1991) J. Amer. Ceram. Soc., 74, 1487. Herman, H. and Bhat, H. (1980) in Synthesis of Metastable Phases, eds. Machlin, E. S. and Rowland, T. J. (TMS- AIME, Warrendale, PA), p. 115. Hirayama, T. (1987) J. Amer. Ceram. Soc., 70, C-122. Jayaram, V. and Levi, C. G. (1989) Acta. Metall., 37, 569. Jones, H. (1973) Rep. Prog. Phys., 36, 1425. Jones, R. T., Barcza, N. A. and Curr, T. R. (1997) Plasma Developments in Africa (Internal Report, Mintek, Randburg, South Africa). Kagawa, M., Honda, E, Onodera, H. and Nagae, T. (1983) Mater Res. Bull., 18, 1081. Knight, R., Smith, R. W. and Lawley, A. (1995) Inter J. Powder Metall., 31, 205. Kong, E C. and Lau, Y. C. (1990) Pure Appl. Chem., 62, 1809. Levi, C. G., Jayaram, V., Valencia, J. J. and Mehrabian, R. (1988) J. Mater Res., 3, 969. Lin, C. K. and Berndt, C. C. (1994) J. Thermal Spray Technol., 3, 75. Mackinnon, I. M. and Reuben, B. G. (1975) J. Electrochem. Soc., 122, 806. Madejski, J. (1976) Inter J. Heat and Mass Trans., 19, 1009. Matejka, D. and Benko, B. (1989) Plasma Spraying of Metallic and Ceramic Materials (John Wiley, New York). McPherson, R. (1973) J. Mater Sci., 8, 851. McPherson, R. (1981) J. Mater Sci., 16, 3141. McRae, L. B., Barcza, N. A. and Curr, T. R. (1985) in Proceedings of the International Conference on Mineral Science and Technology, Vol. 2, ed. Haughton (Council for Mineral Technology, Randburg, South Africa), p. 827. Meetham, G. W. (1988) Mater Design, 9, 308. Metaxas, A. C. (1996) Foundations ofElectroheat: A Unified Approach (John Wiley, Chichester, UK). Meyer, D. E., Dillon, R. O. and Woollam, J. A. (1989) J. Vac. Sci. TechnoI., AT, 2325. Miller, R. A., Smialek, J. L. and Garlick, R. G. (1981) inAdvances in Ceramics, Vol. III, Science and Technology ofZirconia, eds. Huer, A. H. and Hobbs, L. W. (American Ceramic Society, Columbus, OH), p. 241. Mohanty, M. and Smith, R. W. (1995) J. Therm. Spray Technol., 4, 384. Mutasim, Z. Z., Smith, R. W. and Mohanty, M. (1992) in Thermal Spray--International Advances in Coating Technology, ed. Berndt, C. C. (ASM International, Materials Park, OH), p. 1019. Notomi, A. and Hisatome, N. (1996) Pure Appl. Chem., 68, 1101. Ohtake, N., Mashimo, Y. and Yoshikawa, M. (1991) New Diamond Sci. Technol., 173, 178. Patil, D. S., Sreekumar, K. E, Venkatramani, N., Iyer, R. K., Ram Prasad, Koppikar, R. S. and Munim, K. R. (1996) Bull. Mater Sci. (India), 19, 115. Pauloski, L. (1995) The Science and Engineering of Thermal Spray Coatings (John Wiley, New York). Pfender, E. (1994) in Proceedings of the International Symposium on Heat and Mass Transfer, ed. Fauchais, E (Bell House, Turkey), p. 223. Ramachandran, K., Patil, D. S., Venkatramani, N., Biswas, A. R., Venkateswaran, S. and D' cunha, R. (1996) Indian Vac. Soc. Bull., 27, 23. Ramachandran., K., Selvarajan, V., Ananthapadmanabhan, P. V. and Sreekumar, K. P. (1997) Plasma Devices and Operations, 5, 1.
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Rangaswamy, S. and Herman, H. (1986) Advances in Thermal Spraying. (Pergamon Press, Oxford, UK), p. 101. Ranz, W. E. and Marshall, W. R. (1952) Chem. Eng. Prog., 48, 141. Rykalin, N. N. (1976) Pure Appl. Chem., 48, 179. Rykalin, N. N. and Kudinov, V. V. (1976) Pure Appl. Chem., 48, 229. Scott, K. T. and Kingswell, R. (1991) in Advanced Surface Coatings, eds. Rickerby, D. S. and Matthews, A. (Blackie, London). Stecura, S. (1977) Amer. Ceram. Soc. Bull., 56, 1082. Taylor, E R., Manrique, M., Pirzada, S. A. and Abdel-Latif, M. (1995) Plasma Chem. Plasma Proc., 15, 545. Turnbull, A. G. and Wadsley, M. W. (1988) The CSIRO Thermochemistry System, version 5 (IMEC, Australia). Venkatramani, N. (1995) Bull. Mater. Sci. (India), 18, 741. Yoshida, T., Kawasaki, A., Nakagawa, K. and Akashi, K. (1979) J. Mater. Sci. 14, 1624. Zaat, J. H. (1983) Ann. Rev. Mater. Sci., 13, 9. Zoltowski, P. (1968) Rev. Int. Htes. Temp. et. Refract., 6, 65.
Chapter 7 Spray-Forming 7.1. 7.2. 7.3.
Introduction Principles Variations and Distinctions 7.3.1 Variations 7.3.1.1 Crucible: Induction Skull Melting/Spray-Forming 7.3.1.2 Atomizer: Circular vs. Linear Spray-Forming 7.3.1.3 Atomizer: Close-Coupled vs. Free-Fall Spray-Forming 7.3.1.4 Atomization Gas: Reactive Spray-Forming 7.3.1.5 Substrate: Near Net Shape Spray-Forming 7.3.1.6 Spray-Forming and Co-Injection 7.3.2 Nomenclature 7.3.2.1 Spray Atomization 7.3.2.2 Related Processing Techniques 7.4. Applicability 7.5. Non-Equilibrium Phenomena in Spray-Forming 7.5.1 Non-Equilibrium Nature 7.5.1.1 Rapid Solidification in Atomized Droplets 7.5.1.2 Transient Semi-Solid Layer in Deposition Stage 7.5.2 Non-Equilibrium Related Features in the Deposit 7.5.2.1 Metastable Phases 7.5.2.2 Extended Solid Solubility 7.5.2.3 Absence of Macrosegregation/Minimized Microsegregation 7.5.2.4 Refinement 7.5.3 Effects of Non-Equilibrium Features on Mechanical/Physical Properties 7.5.3.1 Direct Effects 7.5.3.2 Indirect Effects 7.6. Concluding Remarks Acknowledgments References
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Chapter 7 Spray-Forming B ING LI and ENRIQUE J. LAVERNIA
7.1. I N T R O D U C T I O N
Materials science and engineering is a very broad, interdisciplinary field which encompasses several sciences, primarily chemistry, physics, and metallurgy. The materials under investigation include metals, semiconductors, ceramics, polymers, glasses, and composites. The areas of study range from the nature of atomic arrangement to the phenomenology of the macroscopic material properties; and from elegant fundamental theories to exciting practical applications. Among the manifold aspects, an everlasting important philosophy in materials science and engineering is to produce a material having the required physical and mechanical properties at the lowest cost. To fulfill this goal, it is essential to explore advanced materials with superior properties while simultaneously maintaining a competitive cost. This is made possible by development/utilization of novel processing techniques. Each of the above mentioned categories of materials has its own unique processing requirements, either conventional or novel. For metallic materials, there are mainly two most frequently utilized conventional processing techniques, i.e., casting (ingot casting, die casting, investment casting, etc.) and powder metallurgy. Meanwhile, many novel processing techniques are being currently explored/investigated, such as mechanical alloying (Koch 1989), laser surface modification (Steen 1985), ion implantation (Picraux 1984), self-sustained high-temperature processing (Mabuchi et al. 1995, Alman et al. 1995, Rawers and Wrzesinski 1992), and spray-forming, to name a few. Each of these novel processing techniques has its own advantages and disadvantages. Among these many techniques, spray-forming is especially attractive in the sense that it is potentially capable of fabricating advanced materials with improved properties at costs that are industrially competitive. The improved properties are directly associated with the non-equilibrium nature of spray-forming, which is the focus of this chapter and will be discussed in detail in the following sections. The cost effectiveness is derived from the inherent simplicity of this technique, which, in turn, is partially attributable to the non-equilibrium phenomena. It is instrumental to illustrate its inherent simplicity by comparing spray-forming with other processing techniques, especially the two well-established processing techniques: powder metallurgy and ingot casting. Compared with powder metallurgy, spray-forming combines powder production, sieving, canning, cold pressing, degassing and hot press153
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Bing Li and E. J. Lavernia
ing into a single step. This is inherent in the concept of spray-forming which incorporates consolidation into the process of powder production. Compared with the conventional ingot casting, spray-forming eliminates macrosegregation and minimizes microsegregation, hence avoiding many intermediate forming steps, such as long term homogenization treatment and extensive thermomechanical processing, which are typically required in conventional ingot casting techniques to eliminate/minimize segregation. The concept of spray-forming was first introduced by Singer in the early 1970s (Singer 1970, 1972). His initial intent was to develop a processing technique which was capable of reducing the overall cost in manufacturing structural materials. It was soon realized that spray-forming is also unique in obtaining superior microstructures and hence improved mechanical properties. Because of the many potential advantages of this technology over conventional processing techniques, spray-forming aroused more and more interest in the academic field, as well as in the industrial community worldwide over the last three decades. There are now more than 20 research organizations and 20 industrial companies located in the United States, Europe, and Eastern Asia (China, Japan, Korea, etc.) actively involved in the development/utilization of this technology (Leatham 1993). These developmental activities are coupled with a large volume of publications in the field of spray-forming, including journal articles, conference papers, book sections (Yule and Dunkley 1994), and books (Lavernia and Wu 1996). These efforts may be categorized as: (a) understanding of the fundamental phenomena associated with this technology, such as heat transfer, undercooling, and solidification; (b) application of this technology to new alloy systems rather than the widely studied aluminum alloys, such as Cu, Fe, Ni, Ti, Zr, and Mo based alloys; (c) optimization of the processing parameters to achieve better microstructural characteristics and mechanical properties; (d) modification of this technology to extend its application domain from monolithic alloys to metal matrix composites; (e) novel design of this technology to enhance its capability in the control of product geometry (near net shape, such as flat sheet, circular tube, etc.); and (f) commercialization of this technology.
7.2. P R I N C I P L E S
The principle of spray-forming is schematically shown in Fig. 7.1. It consists of two stages: spray atomization and spray deposition. Accordingly, spray-forming is also termed as spray atomization and deposition in the academic field as well as the industrial community. The spray atomization stage further contains the following detailed physical changes: (a) A stream of molten metallic material is impinged upon by highly energetic gas jets in different, normally symmetrically opposite directions, as shown in Fig. 7.1. The result is disintegration of the molten metallic stream into small, irregular ligaments. (b) The small, irregular ligaments transform into spherical droplets shortly after their formation (Lavernia et al. 1992). This change is normally termed spheroidization in
Spray-Forming
155
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.
.
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.
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.
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~ ~
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Figure 7.1. Two stages of spray-forming: sprayatomization and spray deposition.
spray-forming. The driving force for this change comes from the decrease of the surface energy associated with the ligament (minimum energy principle). The surface energy is the product of surface tension and surface area. The surface tension is a constant for the same material, while the surface area depends on the shape. For a given volume, a sphere has the minimum surface area. Accordingly, there is always a tendency for any shape to change into spherical to minimize the total surface energy. In other words, formation of spherical droplets from irregular ligaments is thermodynamically favorable. In a practical situation, whether spheroidization really occurs depends on the associated kinetics. It has been demonstrated, both experimentally and theoretically, that spheroidization can be completed almost immediately after the formation of the irregular ligaments during spray-forming, mainly because of the high surface tension of most molten metallic materials and the size of the ligaments. (c) The sizes of the spherical droplets are normally on the order of microns, but typically fall into a wide range. Ample experimental evidence suggest that, in any spray-forming experiment, the size distribution of these droplets obeys a specific distribution called log-normal distribution (Lubanska 1970, Liang et al. 1992a, Marinkovich et al. 1989, Irani and Callis 1963, Allen 1981, Grant et al. 1993a, Delplanque et al. 1997). Most of the droplets have sizes close to a certain finite value (median size). When the size deviates further from this value, the probability of finding a droplet having this size decreases. (d) The atomization gas flow field forms a specific geometrical shape, normally a spray cone. The micron-sized droplets are dispersed in the spray cone. (e) The spherical droplets experience momentum transfer with the atomization gas flow field. This momentum transfer governs the movement of the droplets in the spray
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Bing Li and E. J. Lavernia
cone, including trajectory, flight time, and subsequent deformation behavior during the deposition stage. (f) The spherical droplets experience thermal exchange with the atomization gas flow field. This thermal exchange causes the spherical droplets to cool down and solidify during their flight towards the substrate. The thermal exchange is affected, to a certain extent, by the momentum transfer between the droplet and the gas field. The spray deposition stage involves collection of the droplets in the spray cone onto a water-cooled substrate, forming a dense bulk material that is normally referred to as a deposit: (a) By the time when these droplets reach the substrate, they are in the solid, liquid or semi-solid states. (b) At the beginning of the deposition, the droplets impinge directly on the top surface of the substrate. Once these droplets are collected on the substrate, they act as the substrate themselves for the subsequent impinging droplets. (c) Impingement of the droplets with the substrate leads to extensive modification of the droplets, including their morphology and the associated microstructural features, such as dendrites (see Fig. 7.2), which are developed during solidification in flight. This deformation plays a critical role in the evolution of the microstructure of the deposit. (d) The deposited bulk material experiences thermal exchange with the water-cooled substrate and the surrounding environment, leading to the dissipation of its thermal energy. Simultaneously, during spray deposition, the arrival of droplets from the spray cone onto the substrate continuously incorporates thermal energy into the deposited bulk material. These two effects compete with each other, determining the overall thermal behavior of the deposit. This further affects some aspects of the microstructure, such as grain size. The microstructure in turn determines the mechanical properties. Spray-forming is normally carried out in an environmental chamber. The chamber is evacuated and back-filled with a protective gas such as N2, He, and Ar. The purpose of this arrangement is to enhance the capability to either suppress or exploit the possible reactions, such as oxidation, between the atomization/environmental gas and the metallic materials (the droplets and the deposit). The products of spray-forming include the bulk deposit and powders that are usually termed as oversprayed powders. During the deposition stage, some powders (solidified droplets) cannot be collected into the substrate to form the bulk deposit. Instead, they are carried by the gas flow field into a collecting apparatus (cyclone). These powders are termed as oversprayed powders in spray-forming technology. In commercial applications, the bulk deposit is the main product of practical value, while the oversprayed powders must be minimized in quantity. In scientific research, both the bulk deposit and the oversprayed powders are useful in providing valuable information on the understanding of the fundamentals in spray-forming.
Spray-Forming
Figure 7.2. SEM
157
micrograph showingdendrites formed in spray atomized A1-Mn-LiZr powders (Lavernia et al. 1992).
7.3. VARIATIONS AND D I S T I N C T I O N S
Advances in our understanding of the science and technology of spray-forming have led to the development of novel designs, as well as modifications to the original technology, although the fundamental principle remains the same. A comprehensive discussion should encompass a description of these advanced designs. Since spray-forming is currently a very active field, an exhaustive description of all these variations is difficult. Accordingly, in the following sections, several representative variations will be described. It is also worth noting that there is some confusion in the literature regarding the term sprayforming, as applied to related manufacturing technologies. This issue is addressed in the sections that follow. 7.3.1 Variations
Most of the variations to spray-forming involve modification of one or more of the key components of a spray-forming facility, which include the following: crucible, atomizer, atomization gas, and substrate, as will be described in sequence below.
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7.3.1.1 Crucible: induction skull melting~spray-forming In spray-forming, there is generally a crucible used to melt the master alloy and then to hold the molten metal. During spray-forming of non-reactive alloys with low to intermediate melting temperatures, such as aluminum and magnesium based alloys, ferritic steels, and nickel based intermetallics, the crucible is generally made of graphite or ceramic materials such as zirconia and alumina. For reactive/refractory alloys, these conventional crucibles are no longer usable. In spray-forming, the molten metal is generally heated to a temperature about 50 ~ 100 K higher than its melting temperature and then maintained there for a significant time duration (> 10 min). Such a processing procedure is essential to ensure smooth flow of the molten metal through a delivery tube into the atomizer or, in the worst scenario, to avoid freeze-up of the molten metal at the delivery tube. For reactive alloys, such as Ti (Morris et al. 1992, Li et al. 1996a, Li and Lavernia 1997) and Zr (Li et al. 1997) based alloys, their highly reactive nature in the molten state will lead to significant reaction between the molten alloys and the crucible material during such a long period, making selection/utilization of conventional crucibles extremely difficult if not completely impossible. For refractory alloys, such as Zr (Li et al. 1997), Mo, Nb, V, Ta (Delplanque et al. 1997), and W based alloys, their melting temperatures (>1800~ are normally close to or above those of the commonly used crucible materials (zirconia or graphite). Accordingly, the crucible will be destroyed before the alloys can be melted. To circumvent these problems, the facility has to be significantly modified, and an alternate melting approach, namely skull melting, has to be incorporated. The novelty in the concept of skull melting/spray-forming may be best illustrated by a comparison between the usage of the conventional crucible and the skull crucible (refer to Fig. 7.3). In spray-forming experiments that employ a conventional crucible, the crucible is charged with a master alloy and then heated up using an induction furnace. In the process of heating, the crucible is thermally insulated from its environment, preventing unnecessary loss of thermal energy and enhancing melting efficiency. Consequently, both the crucible and the charge remain at an elevated temperature as they are heated up. In skull melting/spray-forming, the crucible is water-cooled during the entire process of heating of the charged material. When the center of the charge is melted, its exterior regions remain solid because of their intimate contact with the cold crucible. These solid regions of the charged material form a crucible of the shape akin to the water-cooled crucible. This crucible is called a "skull crucible", which leads to the name of this technique. The most intriguing feature of skull melting technique is that the molten metal will be contained in a crucible (the skull crucible) made of the metal itself; hence contamination resulting from the reaction between the molten metal and the crucible is minimized, if not totally eliminated. 7.3.1.2 Atomizer: circular vs. linear spray-forming
In terms of the geometrical shape of the atomizer, there are two typical types of sprayforming techniques: circular and linear. In a circular atomizer, the gas jet nozzles are arranged concentrically (Fig. 7.4a). The number, cross-section area, and shape of the nozzles are critical parameters that control the efficiency of the atomizer. During sprayforming, the atomization gas streams exit these nozzles, forming a circular cone (Fig. 7.5b). The deposit produced this way normally exhibits a bell shape (Fig. 7.5b). In a
Spray-Forming
159
Thermocoupl, Moltenmetal
Stopper Moltenmetal Ceramic cruci Thermalinsul Inductioncoil
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(a)
(b)
Figure 7.3. Comparison on the usage of: (a) conventional ceramic crucible and (b) skull crucible.
linear atomizer, the gas jets are located along two parallel lines (Fig. 7.4b). During sprayforming the atomization gas forms a prism-like geometry, giving rise to a deposit with a high aspect ratio (Fig. 7.5a). Hence, a linear nozzle, in combination with a moving substrate, may be readily used to attain a sheet-like geometry (Fig. 7.5a).
7.3.1.3 Atomizer: close-coupled vs. free-fall spray-forming Schematic diagrams illustrating these two types of spray-forming are shown in Fig. 7.6. In close-coupled spray-forming, the molten metal is released from the crucible through a delivery tube. When it exits the tube, the high-speed atomization gas jets impinge on the melt stream at the tip of the tube (Fig. 7.6a). In free-fall spray-forming, the molten metal is released from the bottom of the crucible, forming a stream. The molten metal stream travels downwards, through the atomizer unit, until it is atomized at a certain point below the atomizer. From the point it exits the bottom of the crucible to the point it is atomized, the molten stream is not constrained by any tubes as in close-coupled atomizer (Fig. 7.6b). Instead, it falls freely, which leads to its name: free-fall sprayforming. Close-coupled spray-forming and free-fall spray-forming are complementary to each other. Close-coupled spray-forming is more susceptible to freeze-up problems. Since the delivery tube is in close contact with the cold atomization gas, there is a higher probability for the melt stream to freeze at the tip of the tube. When this happens, further delivery of the molten metal through the tube is impossible, and the entire experiment fails and has to be restarted. Free-fall spray-forming can avert freeze-up problem because of the absence of the delivery tube. A concomitant problem is the instability of the molten stream in the absence of the delivery tube. The molten metal stream may be easily deflected by the turbulence in the surrounding gas envelope. This instability may lead to oscillation and asymmetry of the atomization cone (Lavernia and Wu 1996).
160
Bing Li and E. J. Lavernia
Gas jet nozzles Orl ac(
delivery tube
(a)
Space to accommodate delivery tube
Linear atomizer twins
~as jet slits
(b) Figure 7.4. Arrangementof atomization gas jet nozzles (slits) in circular and linear atomizers: (a) circular, and (b) linear.
7.3.1.4 Atomization gas: reactive spray-forming There are two types of atomization gases: protective and reactive. Helium and argon are common protective gases, while 02, CO, CO2 and H2 are normally categorized as reactive gases. Nitrogen (N2) is either protective or reactive, depending on the nature of the material being synthesized. In many situations, the atomization gas employed is protective. In a recently developed variation, i.e., reactive spray-forming (Zeng et al. 1992, Zeng et al. 1995), the atomization gas is deliberately selected to contain certain amount of reactive gases. Reactive spray-forming is mainly utilized to synthesize a special category of high strength materials: dispersion-strengthened alloys. The most representative example in this category is the oxide dispersion-strengthened (ODS) material. Its earliest application may be dated back to 1910 when Coolidge incorporated thorium oxide into tungsten wire to improve its elevated temperature mechanical properties (Coolidge 1910). The first
Spray-Forming Molten metal
161
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(a) Molten metal Delivery tube Circular atomizer
Atomized droplets . . . . .
Deposit Water-cooled copper substrate -Rotation and/or translation movement
......
~,._~
Figure 7.5. Linear vs. circular spray-forming: (a) linear (Lavernia and Wu 1996), and (b) circular.
162
Bing Li and E. J. Lavernia
('ruciMc
Liquid metal
Crucit"
~netal
v tube Ammiz assemM
Pro-heated atomizing assembly
~S
Inert gas /
Spray of atomized /~-.-droplets //
Refractory t ubc
Spray dropl,
\
Figure 7.6. Close-coupled (a) and free fall (b) spray-forming (Singer, 1970, Lavernia and Wu 1996).
structural ODS, sintered aluminum powder (SAP), was developed by Irmann in 1952 (Irmann 1952). In 1970, Benjamin introduced the concept of fabricating ODS using mechanical alloying (Benjamin 1970). Since then, many ODS systems, notably Ni- and Fe-based alloys, have been successfully developed. It should be noted that in either of these conventional techniques, typical powder metallurgy processing procedures, such as powder production, canning, degassing, and consolidation, have to be carried out. In reactive spray-forming, the atomization gas reacts with the droplets, forming inorganic phases such as oxides, carbides, or nitrides. The large surface area associated with the micron-sized droplets greatly facilitates the reaction kinetics. Nevertheless, this is offset to some extent by the limited time of reaction during spray-forming, since the cooling and solidification rates of the droplet are very high. For example, a 20 #m A1 droplet may be cooled down from 800~ to 80~ in about 6 ms. Consequently, only a very thin intermetallic film (0.1-30 nm) may be formed on the surface of the droplet. This ensures that the size of the intermetallics is fine (<0.1 #m), and hence presumably beneficial to the mechanical properties. The presence of coarse intermetallic particles generally has detrimental effect on the mechanical properties. During deposition stage, the thin film will be deformed and fragmented due to the high velocity impact of the droplets with the substrate/deposit, leading to a uniform distribution of the intermetallic particles in the deposit. A uniform distribution of the intermetallic particles in the matrix is another criteflon that must be satisfied for them to be beneficial to mechanical properties. The main advantage of reactive spray-forming over powder metallurgy for the fabrication of dispersion strengthened materials lie in its simplicity: it combines powder production, reaction, canning, degassing, and consolidation into a single step.
Spray-Forming
163
Induction
heated ladle
Tube
.--L
~x- ~ -
~.....
AKomizer
_~_
Recipient Deposition chamber
- - - - Exhaust
Figure 7.7. Spray-forming of tubular materials (Lavernia and Wu 1996, Fiedler et al.
1987).
It is worth noting that even in reactive spray-forming, the protective environment in the chamber mentioned in Section 7.2 remains necessary. In the absence of a protective environmental chamber, the ratio of the reactive to the protective gas will change unpredictably. This unpredictability is to be avoided either in scientific research or scale-up commercial production.
7.3.1.5 Substrate: near net shape spray-forming As stated earlier, segregation (macro- and micro-) is inherently minimized in sprayforming, ensuring microstructural homogeneity in the deposit. Such a homogeneity leads to another important feature of spray-forming: its potential capability for net shape or near net shape forming, which means that the products of spray-forming need no, or limited, further thermomechanical treatment. The importance of this feature lies in the associated cost reduction when applied to a commercial product. However, full utilization of this potential depends on proper design of the spray-forming facility, especially the geometrical shape and the movement of the substrate. Representative examples are shown in Fig. 7.5a and Figs. 7.7-7.10. From these designs, flat sheets (Fig. 7.5a), round tubes (Fig. 7.7), strip ribbons (Fig. 7.8), rolls (Fig. 7.9), and round billets (Fig. 7.10) may be manufactured in near net shape. 7.3.1.6 Spray-forming and co-injection This technique was developed during the course of synthesizing metal matrix composites (MMCs) using spray-forming techniques. Metal-matrix composites may be synthesized exactly the same way as that employed to synthesize the monolithic alloys using the conventional spray-forming technique. In this case, the charge in the crucible is either a pre-synthesized metal matrix composite bulk material or a mixture of the monolithic alloy and the reinforcement. Challenges to this method include:
Bing Li and E. J. Lavernia
164
Liquid metal
Crucible
Close-type atomizer
Inert gas Chamber
Atomized droplet spray
Spray-deposited strip Roll
Roll, substrate Hydraulic ram Hot rolled strip
Figure 7.8. Spray-forming of rolled strips (Singer 1970, Singer 1972, Lavernia and Wu 1996).
Chamber
9
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F~refo~~ _~'-----~-
I//.O'/Z~7/////I'~,,,'$.
i (a)
'
"//.////h (b)
Figure 7.9. Spray-forming of mill rolls: (a) schematic diagram of the equipment and (b) rolls produced (Lavernia and Wu 1996, Ikawa et al. 1990).
165
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(c) Figure 7.10. Spray-forming of round billets using (a) horizontal substrate plus vertical motion (Lavernia and Wu 1996), (b) inclined substrate plus vertical motion (Lavernia and Wu 1996, Leatham and Lawley 1993), and (c) inclined substrate plus horizontal motion (Leatham 1996).
i
166
Bing Li and E. J. Lavernia
(a) Long time exposure of the reinforcement to the matrix environment at elevated temperatures, leading to detrimental interfacial reactions between the reinforcement and the matrix. (b) The difference in melting temperatures of the matrix and the reinforcement. In many cases, such as A1/SiC, the difference is so large that it is not practical to melt the matrix and the reinforcement simultaneously. Accordingly, the matrix is melted while the reinforcement remains solid. In the extreme, when the volume fraction of the reinforcement is high, the fluidity of the melt becomes an issue. In this case, it will be very difficult for the melt containing the reinforcement to flow through the delivery tube. Therefore, it is difficult to synthesize MMCs with high volume fraction of reinforcement using this method. As an alternative to circumvent these problems, spray-forming and co-injection was developed in the 1980s (Singer and Ozbek 1985, Willis 1988, Gupta et al. 1989, 1990, Wu and Lavernia 1991). Spray-forming and co-injection does not involve modification of any of the key components of spray-forming mentioned earlier, i.e., crucible, atomizer, atomization gas, and substrate. Instead, it incorporates a new component: the particulate injection unit, as illustrated in a schematic representation of the spray-forming and co-injection facility used to fabricate TiA1/TiB2 composites (Fig. 7.11). The particulate supplying compartment is connected to a high-pressure protective gas on one side, and injector tubes the other side. The injector tubes are positioned at a certain distance below the atomizer. In order to obtain uniform distribution of the reinforcement in the composite, these injector tubes are normally positioned symmetrically around the central line of the spray cone. During the experiment, the crucible is charged with the matrix alloy. Upon melting, the molten matrix alloy is released through the delivery tube and subsequently atomized, forming the spray cone. Following atomization, as the partially solidified matrix droplets travel towards the water-cooled substrate under the combined forces of gravity and fluid drag, the injector tubes release high speed gas flow which contains entrained reinforcement particulates. The reinforcement particles penetrate into the matrix droplets, forming composite droplets. The micron-sized composite droplets experience cooling and solidification, and are finally collected on the substrate as a highly dense composite preform.
7.3.2 Nomenclature 7.3.2.1 Spray atomization The term spray-forming is sometimes confused with spray-atomization in the scientific literature. Possible reasons include: (a) The term spray-atomization has a much longer history than spray-forming. Its earliest utilization may be traced back to the 1920s (Lawley 1992). (b) Spray-forming was originally developed on the basis of spray atomization by incorporating a substrate across the spray cone to collect the impinging solid/semisolid/liquid droplets. Hence, in the absence of deposition stage, spray-forming becomes spray atomization.
167
Spray-Forming En vi ron mental ch a
m
b
_
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Skull crucible ..............{~z].....................i.....................~.. Moltenmelt . ~ ~ I ~ Induction coil (2) ~ ~\ - ) ' ' ~ - ~ ~ ~t
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............. I
~ j , / j ~ ~,"\~ .....~ i
HIPOWDER INJECTOR COLUMN -3"-~t ' / ~] llnert gas cxhau,t ;
.- CycJo.e ! collector CYCLONE SEPARATOR Figure 7.11. Schematic diagram showing the spray-forming and co-injection facility used to fabricate TiA1/TiB2 composites (Li and Lavernia 1997).
(c) In many cases, the exterior appearance of spray-forming facility is the same as that of a spray-atomization facility. However, from a processing point of view, spray-forming is completely different from spray-atomization. Spray-atomization is a powder production technology. It is only a method to synthesize powders which will be used in powder metallurgy technology. It is only one of the many steps in powder metallurgy processing technology. In comparison, spray-forming is a technology that parallels conventional casting and powder metallurgy, with many potential advantages over conventional casting and powder metallurgy. 7.3.2.2 Related processing techniques
Other spray-related processing techniques include: plasma spray-forming (Heimann 1996, Sampath et al. 1991, Zhao et al. 1993a), the high-velocity oxyfuel thermal spray process (HVOF) (Parker and Kutner 1991), and the electric arc spray process (Buhrmaster et al. 1988, Fussell et al. 1994, Zhao et al. 1993b). The primary differences between these techniques and spray-forming are summarized in Table 7.1. The most critical difference is in the dimension of the product of these processing techniques. One of the main reasons that spray-forming has attracted increasing attention is that it is capable of fabricating bulk materials, just like the conventional casting and powder metallurgy processing techniques.
168
Bing Li and E. J. Lavernia
Table 7.1 Differences between various spray related processing techniques
Techniques
Plasma
High velocity oxyfuel
Electric arc
spray-forming
thermal spray process
spray process
Melting method
Plasma arc
Internal combustion jet
Electric arc
Induction furnace*
Feedstock
Powders
Powders
Consumable
Bulk material
Coating
electric wires Coating
Tonnage bulk material
Product
Coating t
Spray-forming
* Sometimes electrical resistance heating. t Sometimes three dimensional, but limited in size.
7.4. A P P L I C A B I L I T Y
In principle, spray-forming is applicable to any material system, especially metallic systems. The only limitation lies in the melting of the material when: (a) the material's melting temperature is extremely high, such as Ta and W; (b) the material reacts with the crucible when it is at an elevated temperature or in the molten state, such as Ti and Zr; (c) the material is not electrically conductive, or the material is not in bulk form (such as powders, wires). In these cases, induction heating may not work properly. These problems may be circumvented by modifying the spray-forming facility, especially the melting unit, such as has been performed in the implementation of the induction skull melting/spray-forming technique (Delplanque et al. 1997, Morris et al. 1992, Li et al. 1996a, Li and Lavernia 1997, Li et al. 1997). Spray-forming has been heretofore employed to manufacture a variety of materials, including A1, Cu, Fe, Mg, Ni, Ti, and Zr based monolithic alloys, Ni-A1, Ti-A1, and Fe-A1 intermetallics, and metal/intermetallic matrix composites. These materials are summarized in Table 7.2 for conventional monolithic alloys and intermetallics, Table 7.3 for novel alloys, and Table 7.4 for composites. The motivations for manufacturing these materials using spray-forming include one or several of the following aspects: (a) refined and uniform microstructure generated by spray-forming which may lead to improved mechanical properties; (b) cost effectiveness of spray-forming; (c) eliminated macrosegregation, minimized microsegregation, and extended solid solubility, which may lead to development of novel alloy compositions, such as in A1-Li, A1-Si, A1-Cu, A1-Zn and Cu-Zr alloy systems; (d) rapid solidification effects in spray-forming which promote improved physical properties, such as electric conductivity of Cu based alloys, magnetic properties in Fe-Nd-B systems, as well as damping capacities in A1/graphite/SiCp and A1/graphite composites; (e) superfine and uniformly distributed oxides in reactive spray-forming, which may be exploited as dispersion strengthening phases;
Spray-Forming Table 7.2
169
Spray-formed conventional commercial/experimental alloys*
Category
Specific alloys
References
A1 based alloys
2014A1 2024A1 2090A1 2618A1 A390A1 5083A1 6061A1 7075A1
Annavarapu and Doherty (1995), White et al. (1988) Lavernia et al. (1986) Lavernia and Grant (1988) White et al. (1988), Underhill et al. (1993) Wu et al. (1995) Wu et al. (1997) Wu and Lavernia (1991), White et al. (1988) Grant and Cantor (1990), Lengsfeld et al. (1995), Chu et al. (1998) Lengsfeld et al. (1995) Lengsfeld et al. (1995) Lengsfeld et al. (1995) Hariprasad et al. (1993) White et aL (1988), White et al. (1994), White et al. (1989) Wu et al. (1992) Ebert et aL (1997) Lavernia et al. (1991), Lavernia et al. (1987) Lavernia et al. (1991), Lavernia et al. (1987) Rickinson et al. (1981) Ikawa et al. (1990) Hanlon et al. (1997) Namba et al. (1993) Brooks et al. (1977a), Asano et al. (1991) Megusar et al. (1984) Rickinson et al. (1981), Ucok et al. (1991) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977a) Ibrahim et al. (1992) Tanigawa et al. (1986) Lawrynowicz and Lavernia (1994) Dalai and Pilchard (1993), Huron (1993), Dalai and Prichard (1992) Dalai and Prichard (1992), Bricknell (1986) Dalai and Pilchard (1992) Dalai and Pilchard (1992), Benz et al. (1993) Dalai and Pilchard (1992) Annavarapu and Doherty (1995) Dalai and Pilchard (1993), Fiedler et al. (1987) Dalai and Pilchard (1992), Ogata et al. (1986) Evans et aL (1985) Chu et aL (1988) Dalai and Pilchard (1993), Dalai and Pilchard (1992) Dalai and Pilchard (1992) Underhill et al. (1995) Underhill et al. (1995) Liang et al. (1992a), Zeng et al. (1995), Vetter et al. (1990), Lawrynowicz et al. (1997a) Li et aL (1997) Morris et al. (1992), Li et al. (1996a) Morris et al. (1992) Delplanque et al. (1997)
Mg based alloys
Fe based alloys
Ni based alloys
7150A1 7090A1 7475A1 8009A1 8090A1 WeldaliteTM 049 QE22 Mg Mg-5.6Zn-0.3Zr Mg-8.4A1-0.2Zr M15 Fe-2.5C-6V Fe-0.8C-(3 ~ 5)Cr Fe-2.3C-18Cr-2Mo-IV Stainless steel: (10 ~ 34)Cr 9Cr-Mo 316L En. 3A En. 8 En. 16 En. 19 En. 24 En. 31 En. 33 En. 58J SAE1020 Others: Fe- 16Nd-8B Fe3A1? R e n g f a m i l y : Ren6 41
H i g h s p e e d tool steel:
Ren6 80 Ren6 88DT Ren6 95 Ren6 125 Inconelfamily:
IN 625
IN 718 Others: Mer176 Nimonic 115 Astroloy Waspaloy AF2 - 1DA-6 UDIMET 720 MAR-M-200 Ni3A1t
Refractory and reactive alloys
Zircaloy-4 TiA1t TiA13t Ta-14.6Fe
* For detailed information on their nominal compositions of these commercial alloys, see corresponding references as well as (Brandes and Brook, 1992). t Intermetallics, in atomic percent.
c
4
0
Table 7.3 Novel commerciaVexperimenta1alloys developed using spray-forming Category
Alloy composition
High strength alloys
N202: AlL(6.1 6.S)Cu-(1.7 1.9)Mn-(0.4 O.S)Mg-(0.9 l.l)Zr(V, Ti) ES2x: AI-(5 6)Cu-(O.S 0.8)Mn-0.5Mg-0.5Ag-l .OZr 1.6)Ni-(1.0 1.2)MgN403: AL(17.5 19)Si-(4.2 4S)Cu-(l. IOFe(Mn, Cr, Ti, Zr, V) AI-20Si-SFe-2Ni AI-20Si-SFe
-
-
References
-
-
-
-
-
-
-
AI-20Si-3.1Cu-1.3Mg-0.3Fe
-
-
-
N707: Al-( 10.8 11.4)Zn-(2.2 2.S)Mg- 1.0 1.2)Cu-(0.25 (O.O-O.Z)Si(Fe) Eural: AlL(10 13)Zn-(S 6.5)Mg-(Cr,Mn, Zr) Eura2: same as Eural, but Cr/Mn content reduced
-
AI-12.4Zn-1.94Mg-1.9SCu-O.lZr Light weight high strength alloy
High strength high electrical conductivity alloys Magnetic alloys
Al-1 lZn-2Mg-lCu-0.3Zr UL30: A1-4Li-lCr-0.7Mg-0.3Zr UL40: Al-4Li-O.2Zr Al-4Li- 1Mg-OSGe-0.2Zr AI-3.5 1Li- 1.08Cu-0.63Mg-0.37Zr Cu-1SNi-SSn Cu-(0.1 0.8)Zr Fe-1SNd-8B Fe-2OCo- 1 SNd-8B
-
-
0.32)Zr-
Report (19904 White et al. (1994) Report (1990b) Zheng and Kahl(l990) Zhou et al. (1991) Estrada and Duszczyk (1990) Report (1990~)
a
3
R
de Sanctis (1991) de Sanctis (1991) Lengsfeld et al. (1995) Machler et al. (1991) White et al. (1994), White et al. (1993) Palmer et al. (1993) Gupta et al. (1990). Chanda et al. (1991) Zhang et al. (1989) Herman and Moms (1994). Hermann et al. (1997) Singh et al. (1991) Tanigawa et al. (1986), Harada et al. (1991) Chin et al. (1986)
.h
171
Spray-Forming
Table 7.4 Metal/intermetallic matrix composites fabricated using spray-forming Matrix
Pure A1
TiA1
Singer and Ozbek (1985), Willis (1988), Gupta et al. (1989), Gupta et al. (1990), Wu and Lavernia ( 1991), Llorca et al. (1994), Ibrahim et al. (1991) Wu et al. (1992), Zhang et aL (1993a), Gupta et aL (1993a), Llorca et al. (1993), Kahl and Leupp (1990), Carbalho et al. (1992) Kim et al. (1994) Gupta et al. (1993b), Gupta et aL (1991a), Gupta et al. (1991b), Gupta et al. (1992a) Wu and Lavernia (1991), Kahl and Leupp (1990), Gupta et al. (1992a), Perez et al. (1992), Zhang et aL (1994), Zhang et al. (1993b), Perez et al. (1993) Wu et al. (1994), Wu and Lavernia (1992), Srivatsan and Lavernia (1994), Gupta et al. (1992), McLelland et al. (1991) Gupta et al. (1992c) Maher et al. (1990) Noguchi et al. (1995) Majagi et al. (1992), Perez and Morris (1994) Bedford et al. (1993) Baskin et al. (1994) Lawley and Apelian (1994), Liang et al. (1992b), Lawrynowicz et al. (1997b) Li and Lavemia (1997)
A1203, graphite, SiC, TiB2, TiC, VC, B4C, WC, Cr3C 2 fiber: steel Mo
Lawrynowicz and Lavernia (1994) Lawrynowicz and Lavernia (1994)
A1-Cu
A1-Fe A1-Li
A1-Mg-Si
A1-Si
A1-Ti A1-Zn-Mg-Cu Mg alloys Cu alloys Fe alloys CoSi Ni3A1
Reinforcement
particulate:
172
Bing Li and E. J. Lavernia
I non-equilibrium [ conditions I 'refined microstructure: grains, secondary phase particles
,
1
absence of macrosegregation minimized microsegregation
T
improved mechanical properties 1-"
[near net shape I I simplicity I ._1potential cost !_. --[ effectiveness ]-Figure 7.12.
Interrelationshipbetween differentfeatures of spray-forming.
(f) short term exposure of reinforcement to elevated temperatures during fabrication of composites using spray-forming and co-injection, which may ameliorate interfacial deterioration; (g) penetration of reinforcement particles into droplets in spray-forming and co-injection technique, which allows extension of the reinforcement particles distribution into the interior of the matrix powders. This enhances the homogeneity of the particle distribution in the matrix. In conventional powder metallurgy techniques, distribution of the particles is limited to exterior of the matrix powders. While motivations (a) and (b) apply to all material systems, motivations (f) and (g) apply exclusively to metal/intermetallic matrix composites. Motivations (c), (d) and (e) are applicable to other specific material systems as described above.
7.5. NON-EQUILIBRIUM PHENOMENA IN SPRAY-FORMING
In terms of the technique itself, the characteristics of spray-forming include simplicity and potential cost effectiveness. In terms of the product, characteristics of spray-forming include refined microstructure, absence of macrosegregation, minimized microsegregation, improved mechanical properties, and near net shape capability. These characteristics are interrelated to each other, as shown schematically in Fig. 7.12, and are all associated with the non-equilibrium conditions of this technique.
Spray-Forming
173
7.5.1 Non-equilibrium nature
7.5.1.1 Rapid solidification in atomized droplets The liquid droplets formed in the spray-atomization stage transform into solid powders as they travel down to the bottom of the chamber. The solidification experienced by the droplets during flight is generally not an equilibrium process. Equilibrium solidification is a process in which the solidification rate is infinitely slow such that complete diffusion in the liquid and in the solid are ensured. Such a process is idealized and could never be realized in practical situations. However, when the solidification rate is relatively slow, such as in ingot casting, the process may be approximated as equilibrium solidification. The solidification rate in the droplets is so rapid that this approximation is no longer valid. Such a rapid solidification rate results from the high cooling effect and the large undercooling which may be experienced by the droplets.
Large undercooling. The very first step of the solidification process is nucleation, i.e., the formation of at least one solid nucleus. (1) The temperature at which the nucleation can be initiated is the nucleation temperature. This temperature is always lower than the liquidus temperature, and the temperature difference between these two is the so-called undercooling (prior nucleation undercooling), A T. It represents the driving force for solidification. In solidification, the solid/liquid interface moves towards the liquid at a velocity (2) U, which may be directly related to the undercooling as follows (Levi and Mehrabian 1982): U =~AT
(7.1)
where ~ is a constant. Eq. (7.1) states that a large undercooling level will lead to rapid solidification. The magnitude of the undercooling level is determined by the physical/thermo properties of the melt, the cooling rate, and most importantly, the nucleation mode (i.e., homogeneous vs. heterogeneous nucleation) (Flemings 1974, Kurz and Fisher 1986). In homogeneous nucleation, the solid nucleus is formed out of the liquid independently; while in heterogeneous nucleation, the solid nucleus is formed by attaching the solid atoms to the surface of heterogeneous nucleants (Flemings 1974, Kurz and Fisher 1986). The heterogeneous nucleants are generally solid materials that come into contact with the liquid, including the container-wall and inclusions in the melt, as is often the case in conventional solidification techniques. According to classic nucleation theory, the undercooling level achievable in homogeneous nucleation is always larger (sometimes significantly larger) than that in heterogeneous nucleation (Flemings 1974, Kurz and Fisher 1986). Consequently, in order to achieve large undercooling levels, homogeneous nucleation mode should be promoted. During spray-atomization, the melt is disintegrated into micron-sized droplets. A large proportion of these droplets undergo solidification during flight. Accordingly, container(l) A nucleus is a solid embryowith size larger than a specific critical value. (2) The interface velocity determines the solidification rate. The solidification rate is defined in the literature as the fraction of solid formed per second. In general, a larger interface velocity corresponds to a larger solidification rate. But the explicit form for their interrelationship depends on the geometryof the solidifying system (e.g., spherical droplet or rectangular block) and the geometryof the interface (planar or dendritic).
174
Bing Li and E. J. Lavernia
wall induced heterogeneous nucleation is avoided during solidification. Moreover, the probability for a droplet to contain heterogeneous inclusions decreases as the droplet size decreases. This may be illustrated by a simple calculation. Suppose the number of inclusions in 1 gram of molten metal is Nt. As this amount of melt is atomized into droplets with different sizes, the average weight of the droplets will be Wa grams, and the total number of droplets may be approximated as 1/ Wa. It is evident that the finer the droplets are, the smaller the average droplet weight will be, hence the larger the total number of droplets. When the droplet size is finer than a critical value, 1/ Wd will be much larger than iV,. Therefore, there will be a large proportion of droplets that are free of inclusions. These droplets are able to experience homogeneous nucleation, attain large levels of undercooling, and hence achieving high interface-velocities/solidification-rates. This concept was first employed by Turnbull and Cech (1950) to investigate undercooling phenomena. Once solidification begins, the heat of fusion is released by the liquid during the phase change reaction. If the energy released cannot be effectively dissipated, it will lead to an increase in the temperature of the melt, a decrease in undercooling, and hence a decrease of the interface-velocity/solidification-rate according to Eq. (7.1). Therefore, to ensure high solidification rate during the entire period of solidification, it is necessary to maintain a high cooling rate as well. High cooling effects. In atomization, the elevated temperature droplet loses thermal energy to the surrounding gas environment via a combination of convection, and radiation. The amount of energy dissipated by convection during a period of dt may be expressed as"
Qlds -- h ( T - To)S dt
(7.2)
where T and To are droplet and gas temperature, respectively; S is the area of the interface between the droplet and the atomization gas, which is mathematically equivalent to the droplet surface area; h is the convection heat transfer coefficient, which may be estimated as follows (Ranz and Marshall 1952): t~ ~fl/2 MI/3
h - ( k g / D ) ( 2 + 0 .... Re ~'Pr )
(7.3)
where NRe(--Opgt)dg/#g) and Npr(=#gCpg/kg) are the Reynolds number and Prandtl number, respectively; kg and Cpg are the thermal conductivity and specific heat capacity of gas, respectively; # g the gas viscosity; D droplet diameter; pg gas density; and Vctg droplet-gas velocity difference. The amount of energy dissipated by radiation during the time period dt is given by:
Q2ds -- 6~
- T~)S dt
(7.4)
where e is the emissivity of the droplet, and cr the Stefan-Boltzmann constant. Other variables have the same meaning as defined earlier. Compared with the convection term (Qls), energy loss by radiation is small when the droplet temperature is not very high (<1000~ Accordingly, in many situations, the radiation term is neglected.
Spray-Forming
175
The consequences of energy dissipation from the droplet to the gas environment are as follows: (a) Cooling of the liquid droplet; A decrease of dT in droplet temperature during the time period dT corresponds to a thermal energy decrease of: Q1
rl -- --Pd VdCpl dT
(7.5)
where Pd is the droplet density, Vd droplet volume, Cpl the specific heat capacity of the liquid droplet. The negative sign is incorporated since the temperature decreases, i.e., dT <0. (b) Solidification of the droplet when it is cooled down to a certain temperature below the liquidus of the droplet; The solidification of a fraction, d f , releases the following amount of energy: Q21 -- H f (pd Vd d f ) (7.6) where H i is the heat of fusion. If there is a solidification temperature range, as is the case of alloy systems, the temperature may decrease as well. Accordingly, the total energy released by the droplet is Q2t rl -- Hf(pdVd d f ) - pdVdCpm dT
(7.6)'
where Cpm is the specific heat capacity of the partially solidified droplet. (c) Further cooling of the solidified droplet following the completion of solidification; The corresponding thermal energy decrease during this period is:
Q3rl - --Pd VdCps d r
(7.7)
where Cps is the specific heat capacity of the solid droplet. Energy conservation requires:
Q1 Q2 ds 4- ds --
I Q~l 2 2t Qrl(Qrl) Qr31
prior to solidification solidification
(7.8)
after solidification
Although it is almost impossible to derive a general solution to Eq. (7.8), an order of magnitude estimate on each of the variables in this equation can lead to an important finding, that the small size of the droplet (D) is the key to achieving a high cooling rate and high solidification rate. For example, during solidification, if the temperature of the droplet remains unchanged (such as eutectic or peritectic solidification) and the radiation effect is neglected, then the energy balance equation becomes:
Q1ds _ Q2rl
(7.9)
df _ 6 kg (T - T0)(2 + 0. t;Arl/2 M1/sa dt - H f pd D D "~'Re ~'Pr J
(7.9) t
or explicitly,
176
Bing Li and E. J. Lavernia
For many metals and processing conditions, H f
,~
3
105 J/kg, 1/2
Pd
~
103 kg/m 3,
1/3
kg ~ 10 -2 W/mK, (T - To) "~ 10 K, and (2 + 0.6Nge Npr ) ranges from 2 to 100. Accordingly, d f / d t is in the range of [(10 -7 ~ 10-5)/D2]s-1, where D is in the units of
m. For spherical droplets, the interface velocity may be roughly estimated as follows: U = D(df/dt)
(7.10)
Therefore, the interface velocity is in the range of [(10 -7 -~ 10-5)/D] m/s. In the case of droplets of about 500/zm, the interface velocity ranges from 0.2 to 20 mm/s. Such a solidification rate may or may not be considered as rapid solidification, depending on the precise definition. For a droplet about 5/zm, the corresponding interface velocity is in the range of 20 ~ 2000 mm/s. Such a solidification rate will undoubtedly fall in the so-called rapid solidification regime. Effects o f rapid solidification. Rapid solidification permits large departures from equilib-
rium, resulting in: (a) formation of non-crystalline phases or metastable crystalline phases; (b) extension of solid solubility beyond the maximum equilibrium value; For example, the maximum equilibrium solid solubility of Cr in A1 is 0.44 at.% at 935 K. Upon rapid solidification, it may be extended to exceed 5 at.% (Jones 1978), and (c) Retention of disordered crystalline structures in normally ordered materials and intermetallic compounds. Rapid solidification may also change the microstructure characteristics, including: (a) modification of phase morphology; (b) refinement of dendrites, dendrite arms, and dendrite arm spacing; (c) refinement of second phase particles (constituents, dispersoids and precipitates). 7.5.1.2 Transient semi-solid layer in deposition stage
As stated earlier, spray-forming includes two stages: spray atomization and spray deposition. The rapid, and by definition non-equilibrium, solidification kinetics that are inherent to the atomization stage were examined in the preceding section. In this section, the non-equilibrium conditions present in the deposition stage will be discussed. During deposition, liquid, solid and partially solidified droplets continuously arrive on the top surface of the deposit/substrate. Identification of the precise mechanisms that govern the formation of the microstructure during deposition has been the subject of intense research activities for many years (Liang et al. 1992a, Estrada and Duszczyk 1990, Singer 1982, Singer and Evans 1983, Annavarapu et al. 1988, Gutierrez et al. 1988, Lavernia 1989, Mathur et al. 1989, Annavarapu et al. 1990, Zhuang et al. 1990, Feng et al. 1992, Leatham and Lawley 1993, Annavarapu and Doherty 1993, Liang and Lavernia 1994). It is now believed that the impinging droplets form a thin semi-solid layer on the top surface of the deposit while the rest part of the deposit underneath the semisolid layer remains solid (see Fig. 7.13). This layer experiences/completes solidification and becomes the main part of the deposit, while newly impinging droplets form another
177
Spray-Forming
gas ~ . . . . . . . ~
,~ Q
"1 II~ I
liquid
!
!
I II I
I dendrite~
II
" I~ I
I
,
~
I
I
I I I
I I I |
I
\ \
l
\
Semi-solid
'ate
Solidified deposit 1
9
deposit T = 1 -
10 K/s
F i g u r e 7.13. Semi-solid layer on the top of the deposit (Liang and Lavernia 1994).
thin semi-solid layer. Such a process continues until no more droplets impinge, i.e., at the end of spray-forming. The main part of the deposit underneath the semi-solid layer experiences cooling in the solid state, with no solidification involved. Experimental measurements and numerical studies indicate that the cooling rate during deposition is in the range from 1 -~ 10 K/s (Liang et al. 1992a, Lavernia and Grant 1988, Gutierrez et al. 1988, Mathur et al. 1989). Such a cooling effect is higher than that in ingot casting (<1 K/s) (Lavernia et al. 1992) but lower relative to that in rapid solidification (103 ~ 106 K/s) (Liang et al. 1992a, Mathur et al. 1989, Li et al. 1996b, Gutierrez et al. 1989, Grant et al. 1993b). Accordingly, the cooling process in the main part of the deposit is normally not considered as non-equilibrium. However, the semi-solid layer represents an inherently non-equilibrium condition in several senses. Firstly, when the high velocity droplets arrive and then join the relatively quiescent top surface of the deposit/substrate, they lose their kinetic energy. A portion of this kinetic energy is absorbed by the semi-solid layer, resulting in violent convection in the layer. Secondly, some of the kinetic energy loss is absorbed by the droplets when they experience extensive deformation. Such deformation may include (Liang et al. 1992a, Lavernia 1989): (1) change on the geometrical shape of the droplets from sphere to disk; and (2) deformation and fragmentation of the dendrites formed in the droplets during the atomization stage.
178
Bing Li and E. J. Lavernia
These phenomena dramatically alter the solidification process which is experienced by the droplets. Thirdly, recent results indicate that each discrete thin semi-solid layer completes its solidification in a relatively short time (10 -1 ~ 10 -3 s) because of its thinness and the effective cooling effect of the environmental gas (Liang and Lavernia 1994). Corresponding interface velocity is estimated to be approximately 1 mm/s. Moreover, the cooling rate experienced by the thin layer is in the range of 10 ~ 105 K/s, depending on the materials being spray formed and the processing conditions (Liang and Lavernia 1994). Such a cooling rate and solidification rate may be considered, to a certain extent, to be non-equilibrium. 7.5.2 Non-equilibrium related features in the deposit
In Section 7.5.1, the non-equilibrium conditions that are present during spray-forming were examined. Manifestations of such a non-equilibrium state in the deposit include: metastable phases, extended solid solubility, absence of macrosegregation/minimized microsegregation, as well as refinement of grains and secondary phase particles; these will be discussed individually below.
7.5.2.1 Metastable phases Since atomized droplets experience rapid solidification during their flight, it is anticipated that some metastable phases will form prior to impingement. During the deposition stage, especially in the main part of the deposit, cooling is not as effective as in the atomization stage, allowing the material to be exposed to high temperatures for a relatively long period. Accordingly, there is a possibility for the metastable phases to be transformed into thermodynamically stable phases. Whether these metastable phases will be retained in the deposit depends on: (i). the tendency of metastable phase formation in the material; and (ii). the kinetics of the metastable phase transformation. The former determines the amount of metastable phases present in the material prior to a metastableto-stable phase transformation. The latter affects the amount of metastable phases that are left in the deposit. Experimentally, metastable phases have been heretofore noted in several spray-formed alloys, including A14Mn and ot'-A13Zr phases in A1-6.5Mn2.35Li-0.65Zr (Ruhr et al. 1990, Baram 1991), ot'-A13Zr phase in 7xxx A1 (Machler et al. 1991), 6-A14FeSi2 phase in A1-20Si-5Fe (Zhou et al. 1991), amorphous/quasicrystalline particles in 2024 A1 (Song et al. 1994), as well as Cu2Sb intermetallic phase in Pb-12Sn- 11.5Sb- 1.25Cu-0.75Ni-0.3Cd-0.2As (Upadhyaya et al. 1997). 7.5.2.2 E x t e n d e d solid solubility
The formation of extended solid solution is another characteristic phenomenon of rapid solidification, which might occur during the atomization stage and then modified during the deposition stage. Whether the extended solid solubility is retained in the deposit or not depends on the relevant precipitation kinetics. Experimentally, it has been reported that many spray-formed alloys exhibit solid solubility levels that are much higher than the corresponding equilibrium values at room temperature (Lavernia et al. 1991, Gupta et al. 1992c, Liang and Lavernia 1994, Ruhr et al. 1990, Baram 199 l, Juarez-Islas et al. 1994). For example, the room temperature equilibrium solid solubility of Zn in Mg is
179
Spray-Forming Table
7.5 Extendedsolid solubilityin spray-formedalloys
Alloy (wt.%)
Solid solubility in deposit (wt.%)
A1-6.5Mn-2.3Li-0.65Zr A1-Ti A1-Fe
2.1 "~ 2.25 (Mn) 0.8 ,~ 1.1 (Ti) 2.2 (Fe)
Maximumequilibrium References solid solubility (wt.%) 1.82 (Mn) 1.00 (Ti) 0.025 (Fe)
Ruhr et al. (1990), Baram (1991) Gupta et al. (1992c) Liang and Lavernia (1994)
less than 1.0 wt.% (Lavernia et al. 1991). In a study on Mg-5.6wt.%Zn-0.3wt.%Zr alloys, the room temperature solid solubility of Zn in Mg in the spray-formed Mg-Zn-Zr deposit was found to be 5.3 wt.% (Lavernia et al. 1991), considerably higher than the equilibrium value. However, in this study, the solubility (5.3 wt.%) is still lower than the maximum equilibrium solid solubility of Mg-Zn alloy system, which is 6.2 wt.% at the eutectic temperature (340~ (Lavernia et al. 1991). In extreme cases, the solid solubility in the deposit may be extended beyond the maximum equilibrium solid solubility, as summarized in Table 7.5. Such an extended solid solubility is comparable to that normally observed in relatively more conventional rapid solidification techniques, such as spray-atomization and melt-spinning (Jones 1978). However, compared with these techniques, the extension of solid solubility in spray-forming is generally less pronounced. For example, in the study of A1-6.5Mn-2.3Li-0.65Zr alloy, while the solid solubility of Mn in the deposit is approximately 2.1 ~ 2.25 wt.% (Ruhr et al. 1990), the corresponding value under similar processing conditions in the atomized powders is in the range of 4.25 ~ 6.00 wt.% (Ruhr and Baram 1991).
7.5.2.3 Absence of macrosegregation/minimized microsegregation While the presence of metastable phases and the extension of solid solubility in spray formed alloys are material dependent, the absence of macrosegregation and minimized microsegregation are inherent to spray-forming technology. These characteristics are achieved through a combined effect of the non-equilibrium conditions that are present in the atomization and the deposition stages. During spray atomization, the droplet is small and the solidification rate is high. Accordingly, segregation, such as that associated with the formation of dendrites, is limited to a micron size scale. Fragmentation of dendrites during the deposition stage minimizes microsegregation. It also plays an important role in multiple nucleation in the semi-solid layer. Multiple nucleation is important in preventing the appearance of macrosegregation during the deposition stage. Finally, the relatively slow cooling rate experienced by the solidified deposit underneath the semi-solid layer exposes the material to an elevated temperature for a relatively long period, allowing certain homogenization of the microstructure of the deposit. Such a homogenization effect further minimizes the microsegregation present in the deposited materials.
180
Bing Li and E. J. Lavernia
Figure 7.14. Optical micrograph showing fine equiaxed grains in spray-formed 7075A1 (Lavernia 1989).
7.5.2.4 Refinement Equiaxed grains. A salient feature revealed by numerous studies is the presence of fine, equiaxed grains (see Fig. 7.14) in the deposit, as summarized in Table 7.6. Again, the presence of fine, equiaxed grains in spray-formed materials is directly related to the nonequilibrium conditions present both in the atomization and in the deposition stages. As mentioned earlier, extensive efforts have been dedicated to the understanding of the evolution of such fine, equiaxed grains (Liang et al. 1992a, Annavarapu et al. 1988, Lavernia 1989, Grant et al. 1991, Grant 1995). While there is some controversy on the details of the fundamental mechanisms, it is widely agreed that formation of fine dendrites during atomization stage, and fragmentation of these dendrites during deposition stage, play key roles in the formation of the fine, equiaxed grains. The formation of fine dendrites in the droplets is mainly attributable to the large undercooling, the high cooling effect, and the rapid solidification present during atomization; all of them represent nonequilibrium conditions. Fragmentation of these fine dendrites, on the other hand, is itself a non-equilibrium phenomenon, as discussed in Section 7.5.1.2. Secondary phase particles. In conventional ingot casting, solute segregation is coupled with a slow solidification/cooling rate, leading to the formation of coarse secondary phase particles. In spray-forming, solute segregation is minimized; the rapid solidification experienced by the droplet tends to refine the microstructural characteristics; the high cooling rate in the semi-solid layer and the intermediate cooling rate in the solidified deposit limits the time for secondary particles to precipitate and grow. All of these characteristics of spray-forming lead to refinement/suppression of coarse secondary phase particles, as
Spray-Forming
181
Table 7.6 Grain size in spray-formed materials Alloy (wt.%) Sn-38Pb Mg-8.4A1-0.2Zr Mg-5.6Zn-0.3Zr A1-4Cu A1-4Si 2014A1 2618A1 5083A1 6061A1 A390 Fe-15Nd-8B Cu-6Ti IN625 IN718, Rene 95 Rene 80 316L stainless steel 316L stainless steel 0.18% plain carbon steel SAE 1020 steel Ni3AI* TiAI* Zircaloy-4
Grain size (/zm) 4,-~10 20~50 50~70 75 45 17 10~25 15 24~44 12.5 0.1~5 46 15 30 10~25 10~40 100 10,-~40 10~50 11 30~60 175t
References Lavernia and Wu (1996), Bewlay and Cantor (1991) Lavernia and Wu (1996), Lavernia et al. (1991) Lavernia and Wu (1996), Lavernia et al. (1991) Lavernia and Wu (1996), Grant et al. (1991) Lavernia and Wu (1996), Wu et al. (1993) Annavarapu and Doherty (1995) Lavernia and Wu (1996), Underhill et al. (1993) Wu et al. (1997) Wu and Lavernia (1991) Wu et al. (1995) Veistinen et al. (1987) Lavernia and Wu (1996), Annavarapu and Doherty (1993) Annavarapu and Doherty (1995) Lavernia and Wu (1996), Fiedler et al. (1987) Lavernia and Wu (1996), Bricknell (1986) Lavernia and Wu (1996), Ucok et al. (1991) Lavernia and Wu (1996), Lawley and Cantor (1986) lbrahim et al. (1993) Ibrahim et al. (1992) Lavernia and Wu (1996), Liang et al. (1992a) Morris et al. (1992), Li et al. (1996a) Li et al. (1997)
* In atomic percent. t Grain size of ingot counterpart is larger than 1000/zm.
well as concomitant modification to their morphology. This has been verified experimentally in several studies, as summarized in Table 7.7. Meanwhile, many studies reported the presence of superfine secondary intermetallic particles in spray-formed materials (see Table 7.8). The refined characteristics of these phases are beneficial to the mechanical properties of the material.
7.5.3 Effects of non-equilibrium features on mechanical~physical properties In the development of spray forming, extensive efforts have been devoted to the evaluation of properties of spray-formed materials, including yield strength, ultimate strength, tensile elongation, elastic modulus, fracture toughness, creep resistance, fatigue strength, electrical conductivity, magnetic properties, and damping capacity. The impetus behind these efforts is the belief that the non-equilibrium related features may have a significant effect on the physical/mechanical properties of spray-formed materials, either directly or indirectly. Direct effects are the difference in mechanical properties resulting from non-equilibrium features alone, with no modifications to the material's chemical composition. Indirect effects involve the difference in mechanical properties resulting from
Bing Li and E. J. Lavernia
182
Table 7.7 Effect of spray-forming on the secondary phase particles Alloy
Particle
Processing*
Morphology
Size (/zm)
A390
Primary Si
IM
Block
30~ 100
Eutectic Si
SF IM SF
Particulate Flake-like elm. ~t
5 Coarse elm.
IM SF IM SF IM
Needle elm. Faceted Non-faceted Needle
Coarse elm. 50~ 100 3~ 1 10"~30
SF
elm.
elm.
IM
Particulate
15~20
SF
Particulate
5 ~6
7075A1+ 1Ni+0.8Zr
A13Zr/A13Ni
Babbitt alloyt High speed tool steel
SnSb MC2, M6C
MC
References Lavernia and Wu (1996), Wu et al. (1995) Wu et al. (1995) Wu et al. (1995) Wu et al. (1995) Lavernia and Grant (1986) Lavernia and Grant (1986) Upadhyaya et al. (1997) Upadhyaya et al. (1997) Ikawa et al. (1990), Leatham et al. (1995), Leatham et al. (1985) Ikawa et al. (1990), Leatham et al. (1995), Leatham et al. (1985) Ikawa et al. (1990), Leatham et al. (1995), Leatham et al. (1985) Ikawa et al. (1990), Leatham et al. (1995), Leatham et al. (1985)
* SF: spray-forming; IM: ingot metallurgy. t Pb- 12Sn- 11.5Sb- 1.25Cu-0.75Ni-0.3Cd-0.2As. $ elm.: eliminated.
Table 7.8 Superfine secondary intermetallic particles in deposited materials* Alloy (wt.%) 2618A1 2618+(0.5 -v 1.0)(Fe, Ni) A1-11Zn-2Mg- 1Cu-0.3Zr Fe-0.8C-3Cr 7xxx
Particle t
Size
A19NiFe AI9NiFe AI7Cu4Ni A13Zr Fe3C, Cr3C r/-MgZn2
--~3/zm 0.5--~3.0/xm 0.5,~3.0/zm 25~45 nm 134 nm ~0.25/xm
* All in as-spray-formed condition. t The particle morphology is normally equiaxed.
References Llorca et al. (1993) Underhill et al. (1993) Underhill et al. (1993) Machler et al. (1991) Hanlon et al. (1997) Lengsfeld et al. (1995)
Spray-Forming
183
Table 7.9 Effect of spray-forming on tensile mechanical properties Alloy
Proces sing*
cry S
crU T S
(MPa)
(MPa)
Elongation References KIC (MPaml/2) (%)
2024A1
SF+T4
402~417 562~586
15~16
31
2024A1 Weldalite TM 049 Weldalite TM 049
IM+T4 SF+T6 IM+T6
310~329 441~469 638 665 500 540
19~20 12 16
31
A390A1 A390A1
SF+Extr.+T6 IM+Extr. +T6
506 374
532 387
3.0 1.3
6061 A1 6061A1
SF+T6 IM+T6
313 280
352 310
15.2 13
Cu-0.1Zr Cu-0.1Zr
SF+TMP IM+TMP
441 390
462 450
6 12
Lavernia et al. (1986) Lavernia et al. (1986) Wu et al. (1992) Wu et al. (1992) Wu et al. (1995) Wu et al. (1995) Wu and Lavernia (1991) Brandes and Brook (1992) Singh et al. (1991) Singh et al. (1991)
* SF: spray-forming; IM: ingot metallurgy; Extr." extrusion; T6: solution heat treated and then artificially aged; T4: solution treated and aged naturally to a substantially stable condition; TMP: thermo-mechanical processed.
modification to the chemical composition of the material, which is made possible by the non-equilibrium conditions present.
7.5.3.1 Direct effects It has been experimentally noticed that in the case of numerous alloy compositions that spray-formed materials exhibit better mechanical properties when compared to those of the ingot metallurgy counterparts. Representative results of the tensile properties are summarized in Table 7.9. It is evident that the yield strength and the ultimate strength values are significantly increased upon spray-forming, while the ductility is either improved or degraded. The latter situation may be acceptable if the loss of ductility does not lead to brittle behavior (<2%). Such an improvement in mechanical properties may be associated with the refinement in grains and secondary phase particles, as well as the absence of macrosegregation and the minimized microsegregation. Refined equiaxed grains formed during spray-forming may lead to improvement of the tensile properties, especially the yield (ultimate) strength, because of a decrease in grain size. The well-known Hall-Petch relationship predicts: ffs - - O'so -k- X D g
n
(7.11)
where as is the strength; as0 the internal stress; X and n are experimental constants; and Dg is the grain size. Since the grain size of spray-formed materials is considerably smaller than that present in ingot materials, there is a potential for the strength of the deposit to be improved. A reduction in grain size also decreases the dislocation pile-up length and consequently lowers the stress concentration at grain boundaries and grain boundary triple junctions, hence promoting homogeneous deformation and preventing early crack nucleation and premature failure due to strain localization effects (Perlinde and Lutjering 1982, Lavernia et al. 1990).
184
Bing Li and E. J. Lavernia
Refinement of the secondary phase particles (constituents, dispersoids and precipitates) also directly influences the mechanical properties. Depending on the degree of coherency, there are two types of secondary phase particles: coherent and incoherent. The strengthening mechanisms of these two types of particles are different. For incoherent particles, moving matrix dislocations cannot penetrate these particles. Accordingly, these particles will not be sheared along with the matrix. Instead, the moving dislocation is forced by the applied stress to bend around the particle and bypass it through the Orowan looping mechanism. The critical resolved shear stress due to this mechanism is inversely proportional to the particle size (Reppich 1993): 750R -- C1 (Gmb) v/ f p (7.12) r where C1 is a dimensionless constant, Gm the shear modulus of the matrix, b the Burgers vector, fp the volume fraction of particles, and r the particle radius. Therefore, refinement of secondary phase particles will lead to strengthening effect on the matrix for incoherent particles. For coherent particles, the moving dislocation can cut through them, i.e., they can be sheared along with the matrix. The critical resolved shear stress due to the cutting mechanism is given by (Reppich 1993): h3/2 P 2v/fpr ~c -- C2 ~/--~b
(7.13)
where C2 is another dimensionless constant, h p is a hardening parameter, and other variables have the same meaning as defined earlier. According to Eq. (7.13), coarse particles will have large strengthening effect. This is true only when the particle size is small. When the particle size becomes excessively coarse, Eq. (7.13) will break down for the following two reasons: (1) As the particle size is increased to a certain point, the coherent particle will become incoherent (Reed-Hill 1973). Normally, the lattice parameter of a secondary phase particle is different from that of the matrix. The interface is coherent if the crystallographic planes of the matrix may be epitaxially extended to the particle. The lattice mismatch causes certain strain on the crystallographic planes near the interface. However, as the particle size increases, the lattice parameter mismatch may no longer be accommodated by the strain. Accordingly, the crystallographic plane of the matrix will be terminated at the interface, i.e., the interface becomes incoherent. Under such a circumstance, looping mechanism will be operative. (2) For coherent particles, both cutting and looping mechanisms may be operative, depending on the required stress level (Reppich 1993). When the particle is very small, the stress predicted by Eq. (7.12) may be higher than that by Eq. (7.13). Accordingly, the cutting mechanism requires a lower stress level, and is operative. As the particle size increases, the stress required for the looping mechanism decreases, while that for the cutting mechanism increases. At a certain particle size, the stress level required by the looping mechanism will be less than that by the cutting mechanism. In this case, even though the particle is still coherent, the looping mechanism will be operative.
Spray-Forming
185
When looping mechanism becomes operative, the strengthening effect decreases with increasing particle size. Therefore, for coherent particles, there is a critical particle size at which the strengthening effect is at maximum level. For the classic precipitation hardening, this critical particle size corresponds to peak ageing condition (Reed-Hill 1973). In practice, this critical particle size is in the range of 0.1 ~ 10 #m. Such a fine size is exactly what may be achieved through spray-forming technology (see Table 7.7). In slow cooling/solidification, coarse secondary phase particles are generally segregated to the grain boundary region (Lavernia et al. 1990). The brittle nature associated with many of these secondary phase particles embrittles the grain boundary, making it extremely susceptible to crack nucleation. Such brittle grain boundaries favor the low energy intergranular fracture over high-energy transgranular dimple rupture, resulting in low fracture toughness and poor ductility (Lavernia et al. 1990). Meanwhile, such segregation is sometimes accompanied by a precipitate free zone (PFZ) near the grain boundary (Lavernia et al. 1990). The PFZ region has an inherently lower strength, leading to abrupt variation in local mechanical properties. It is also clear from the above discussion that the compositional uniformity (i.e., absence of macrosegregation and minimized microsegregation) in the deposited materials are critical to the mechanical properties as well. These characteristics decrease the chance of localization of flow stresses, leading to improvement on the mechanical properties, especially the crack resistance, the fracture toughness, and the ductility.
7.5.3.2 Indirect effects
In many alloy systems, the mechanical properties may be improved through effectively modifying the content and the type of alloying elements, provided that the alloying elements are uniformly distributed in the matrix. In ingot metallurgy, such alloying modification is hindered by the macrosegregation of solutes and by the formation of coarse primary or secondary phases. Especially, increased alloying is typically precluded. Sprayforming inherently avoids macrosegregation and microstructure coarsening. Moreover, it sometimes leads to extended solid solution in the deposited materials. Accordingly, sprayforming is an effective method in alloying modifications, especially when heavy alloying is involved. Representative examples in this category include A1-Cu, A1-Zn and Cu-Zr alloys. Commercial A1-Cu alloys (2xxx series) contain Cu, Mn, Cr, Si, and other elements from trace to significant amounts. Further increase in the alloying amounts may improve the available mechanical properties of this alloy system. Nevertheless, exploration/utilization of these potential benefits are difficult in ingot metallurgy because of strong solute segregation. In contrast, compositional uniformity in spray-formed materials enhances alloying flexibility, leading to the development of several heavily alloyed A1-Cu alloys, such as N202 alloy {Al-(6.1 ~ 6.5)Cu-(1.7 ~ 1.9)Mn-(0.4 ~ 0.5)Mg-(0.9 1.1)Zr(V, Ti)} and the family of ES2x alloys {Al-(5 ~ 6)Cu-(0.5 ~ 0.8)Mn-0.5Mg0.5Ag-I.0Zr(V, Ti)}. Comparisons of the mechanical properties of the commercial and novel A1-Cu alloys are provided in Table 7.10, Fig. 7.15 and Fig. 7.16. It is evident there that the mechanical properties of A1-Cu alloys were improved upon heavy alloying.
186
Bing
L i a n d E . J. L a v e r n i a
Table 7.10 Tensile properties of commercial and novel A1-Cu alloys Alloy N202 2618 A1 2014 Alt 2024 A1 2618 A1
Processing*
ffYS (MPa)
ffUTS (MPa)
SF+T6 SF+T651 IM+T6 IM+T6 IM+T6
490 418 385~465 395 330
560 445 440,~500 475 430
References Report (1990a) Llorca et al. (1993) Brandes and Brook (1992) Brandes and Brook (1992) Brandes and Brook (1992)
* SF: spray-forming; IM: ingot metallurgy; T6: solution heat treated and then artificially aged. t Depending on the geometry of the product: sheets, bars, rivet stocks, tubes or wires. 500
400
:--.-.7.---'-~ . . . . . . . . . . . . . . x~ "', x~k
-
"
t~
..........
----,
....
2014
AI
2 6 , s A~
ES20 ES21
300 G
.u ~"
100
0
Figure 7.15.
50
100
150 200 250 Temperature (*C)
300
350
Yield strength of ES2x alloy following thermal exposure at testing temperatures for 100 h (White et al. 1994).
400
380 ... 360 ~" 320 300
.~ 2 8 0
~
-
o'.
-
O
260 -
,
ES20. 176~
240
o
2618 AI. 170"C
220-
9
2618 AI. 176~
200 O.l
J
I
O
J
10
J
100
1000
Time to Rupture (hr)
Figure 7.16. Creep rupture strength of ES20 compared with that of 2618 A1 (White et al. 1994).
187
Spray-Forming
Table 7.11 Tensileproperties of commercial and high Zn content 7xxx alloys Alloy
N707 Eurall Eural2 A1-11Zn-2Mg-1Cu-0.3Zr 7075A1
Processing*
SF+T6 SF+T6 SF+T6 SF+T6 IM+T6
cryS
crUT S
Elongation
(MPa)
(MPa)
(%)
760 775 762"~790 798~810 807 819 705 719 583 600
8.0 2.2~4.9 7.1 29 9
References
Report (1990c) de Sanctis (1991) de Sanctis (1991) Machleret al. (1991) Domalavage et al. (1975)
High strength A1-Zn alloys (7xxx series) contain a significant amount of Zn, Mg and Cu which form major strengthening particles such as intermediate r/t and equilibrium r/phases. Further improvement on the tensile properties may be achieved through either refining grain size or increasing the amount of fine precipitates in the matrix. The latter involves increase in the Zn content to the theoretical limit as predicted in the equilibrium phase diagram while simultaneously maintaining proper ratios of Cu and Mg. In conventional ingot metallurgy, the Zn content is normally below 6.2 wt.%. An increase beyond this limit is difficult because of extensive macrosegregation and cracking associated with slow solidification. These problems, however, may be circumvented by spray-forming. Accordingly, several high Zn content 7xxx alloys have been developed using spray-forming, including N707 {AI-(10.8 ~ 11.4)Zn-(2.2 ~ 2.5)Mg-(1.0 ~ 1.2)Cu(0.25 ~ 0.32)Zr-(0.0-0.2)Si(Fe) }, Eural/Eura2 (3) {AI-(10 ~ 13)Zn-(5 --~ 6.5)Mg-(Cr, Mn, Zr) }, and AI-11Zn-2Mg-lCu-0.3Zr. Comparison of the mechanical properties between the commercially available A1-Zn alloys and the high Zn content alloys is provided in Table 7.11. As shown in Table 7.11, a significant improvement in tensile properties was achieved through an increase of Zn content using spray-forming. Cu-Zr alloys are of interest for applications that require high strength and high electrical conductivity. The high strength of this alloy comes from age hardening effects by fine, zirconium-rich precipitates. However, the equilibrium solid solubility of Zr in Cu is relatively low, with the maximum value being 0.124 wt.% at 950~ When the fine zirconium rich precipitates are produced through solution treatment and subsequent ageing, their amount is restricted by such a small solid solubility. Consequently, the achievable strength is limited. This may be further explained as follows. During ingot casting, solute segregation and slow cooling rate lead to the formation of coarse Zrrich particles in the matrix. Such coarse particles are detrimental, instead of beneficial, to the mechanical properties (see Section 7.5.3.1). When the Zr content is low, all of these coarse particles may be dissolved upon solution treatment. Upon ageing, fine particles are precipitated, enhancing the strength of the alloy. However, when the Zr content is beyond the maximum equilibrium solid solubility (e.g., 0.5 wt.% Zr), only a small amount of the coarse particles may be dissolved upon solution treatment, while others (3) The compositions for Eural and Eura2 are similar, but the content of Cr and Mn in Eura2 is less than in Eural.
188
Bing Li and E. J. Lavernia
Table 7.12 Tensile properties of commercial and high Zr content copper alloys (Singh et al. 1991) Alloy (wt.%) Cu-0.8Zr Cu-0.4Zr Cu-0.2Zr Cu-0.1Zr Cu-0.1Zr
Processingt
crys (MPa)
crUTS (MPa)
Elongation (%)
Conductivity (%IACS)*
SF+TMP SF+TMP SF+TMP SF+TMP IM+TMP
511 525 476 441 390
538 552 504 462 450
3 7 8 6 12
75 82 78 88 85
* IACS: Intemational Annealed Copper Standard. t SF: spray-forming; IM: ingot metallurgy; TMP: thermal mechanical processing, including cold rolling (70%), recrystallization (930 ~ cold rolling (90%), and ageing (400~615 ~
remain in the solvent matrix. Upon subsequent ageing, the remaining coarse particles act as pre-existent nuclei, making it unlikely to generate fine particles. Even though some fine particles are precipitated upon subsequent ageing, their beneficial effect is negated by the detrimental effect of the coarse particles, with the overall effect being hardly beneficial. Such problems inherent in conventional ingot casting may be circumvented through several non-equilibrium features present in spray-forming: extended solid solubility, eliminated macrosegregation/minimized microsegregation, and refined secondary phase particles. Table 7.12 summarizes the high Zr content Cu-Zr alloys produced heretofore by spray-forming, and their significantly improved tensile properties. From the preceding discussion, it is evident that the non-equilibrium characteristics of spray-forming have many direct or indirect beneficial effects on the properties of a material. It is to be noted that many of these features are also present in materials synthesized using other relatively conventional rapid solidification techniques, such as spray atomization and melt spinning. As a matter of fact, some of them such as refinement and extended solid solubility are more pronounced in these techniques than in spray forming because of the higher cooling rates associated with them. Consequently, it may be argued that products of another structural processing technique, powder metallurgy, will exhibit better physical/mechanical properties than spray-formed counterparts, since in many cases the powders used in this technique are produced using spray atomization. This may or may not be true, depending on many factors such as material chemistry and consolidation details. It is a challenge to ensure that the non-equilibrium features obtained in spray atomization are not eliminated in subsequent consolidation steps which normally involve long periods of exposure of the powders to elevated temperatures. Moreover, it is worth noting that, besides the high cost of powder metallurgy, oxidation and contamination (such as pick up of coarse ceramic inclusions during screening (Kim 1988)) problems in this technique may also prevent full attainment of the beneficial effects offered by the non-equilibrium conditions. Oxides present in a material may strengthen the material if they are superfine and uniformly distributed, as in ODS materials (see Section 7.3.1.4). However, the oxide films formed on the surface of powders in conventional powder metallurgy technology often leads to less than optimal fracture properties, if not properly dispersed (Kim 1988, Kim et al. 1985).
Spray-Forming 7.6. C O N C L U D I N G
189
REMARKS
Spray-forming is a novel but powerful processing technology suitable for the manufacturing of structural materials. It has many merits over well established processing techniques such as ingot casting and powder metallurgy. These merits include superior microstructures, improved mechanical properties, and cost effectiveness. Noticeably, many of these merits are associated with the non-equilibrium phenomena present in this technology. Compared with many other novel processing techniques which are being currently explored/utilized, the uniqueness of spray-forming is its capability of manufacturing structural materials in tonnage production levels. This enables it the potential to challenge parallel conventional processing techniques such as casting and powder metallurgy, arousing intensive interest in the industrial community. Spray-forming is also of great value to scientific research community. The rapid solidification effect experienced by the atomized droplets, the transient s e m i - s o l i d region on the top surface of the deposit, the spray-forming and co-injection of m e t a l / i n t e r m e t a l l i c m a t r i x composites, and the reactive spray-forming of d i s p e r s o i d - s t r e n g t h e n e d alloys, all represent areas rich with fundamental problems of interest to the scientific community. Over the last three decades, spray forming has evolved from a laboratory curiosity to a potentially revolutionary manufacturing technique. It is also evident, however, that our understanding of the relevant fundamental physical phenomena involved is, in many cases, inadequate, partly because of the complex interactions of the fluid, thermal, and solidification phenomena that are present during spray-forming. Fortunately, this has not hindered the application of spray-forming technology, as evidenced by the development of a variety of materials with unusual combination of microstructure and physical attributes.
ACKNOWLEDGMENTS
This work was financially supported by the Army Research Office (Grants No. DAAG5598-1-0088 and DAAG55-98-1-0141) and the National Science Foundation (CTS 9614653).
REFERENCES
Allen, T. (1981) Particle Size Measurement (Chapman and Hall, London, UK). Alman, D. E., Dogan, C. E, Hawk, J. A. and Rawers, J. C. (1995) Mater. Sci. Eng., 192/193A, 624. Annavarapu, S. and Doherty, R. (1993) Internat. J. Powder Metall., 29, 331. Annavarapu, S. and Doherty, R. D. (1995) Acta Metall. Mater., 43, 3207. Annavarapu, S., Apelian, D. and Lawley, A. (1988) Metall. Trans., 19A, 3077. Annavarapu, S., Apelian, D. and Lawley, A. (1990) Metall. Trans., 21A, 3237. Asano, I., Araki, T. and Ikawa, Y. (1991) Mater. Sci. Eng., 140A, 461. Baram, J. (1991) Metall. Trans., 22A, 2515. Baskin, D., Wolfenstine, J. and Lavernia, E. J. (1994) J. Mater. Res., 9, 362. Bedford, G. M., Sharp, R. E, Wilson, B. J. and Elias, L. G. (1993) in Spray forming 2, ed. Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 309. Benjamin, J. S. (1970) Metall. Trans., 1, 2943. Benz, M. G., Sawyer, T. E, Carter, W. T., Zabala, R. J. and Dupree, P. L. (1993) in Spray Forming 2, ed. Wood, J. V. (WoodheadPubl., Cambridge, UK), p. 171. Bewlay, B. E and Cantor, B. (1991) J. Mater. Res., 6, 1433.
190
B i n g L i a n d E. J. L a v e r n i a
Brandes, E. A. and Brook, G. B. (1992) Smithells Metals Reference Book (7th edition, Butterworths, London, UK). Bricknell, R. H. (1986) Metall. Trans., 17A, 583. Brooks, R. O., Moore, C., Leatham, A. O. and Coombs, J. S. (1977a) Powder Metall., 2, 100. Brooks, R. G., Leatham, A. G., Coombs, J. S. and Moore, C. (1977b) Metall. Met. Forming, 44, 157. Buhrmaster, C. L., Clark, D. E. and Smartt, H. B. (1988) YOM, 40(11), 44. Carvalho, M. H., Marcelo, T. and Carvalhinhos, H. (1992) Y. Mater. Sci., 27, 2101. Chanda, T., Lavernia, E. J. and Wolfenstine, J. (1991) Scripta Metall. Mater., 25, 131. Chin, T. S., Hara, Y., Lavernia, E. J., O'Handley, R. C. and Grant, N. J. (1986) Y. AppI. Phys., 59, 1297. Chu, M. G., Denzer, D. K., Chakrabarti, A. K. and Billman, E (1988) Mater. Sci. Eng., 98A, 227. Coolidge, W. D. (1910) Proc. Am. Inst. Elec. Eng., 961. Dalal, R. P. and Pilchard, P. D. (1992) in Superalloys 1992, eds. Antolovich, S. D., Stusrud, R. W., MacKay, R. A., Anton, D. L., Khan, T., Kissinger, R. D. and Klarstrom, D. L. (TMS, Warrendale, PA), p. 205. Dalal, R. P. and Pilchard, P. D. (1993) in Spray Forming 2, ed. Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 141. de Sanctis, M. (1991) Mater. Sci. Eng., 141A, 103. Delplanque, J. P., Cai, W. D., Rangel, R. H. and Lavernia, E. J. (1997) Acta Mater., 45, 5233. Domalavage, P. K., Grant, N. J. and Gefen, Y. (1975) Metall. Trans., 14A, 1599. Ebert, T., Moll, E and Kainer, K. U. (1997) Powder Melall., 40, 126. Estrada, J. L. and Duszczyk, J. (1990) Y. Mater. Sci., 25, 1381. Evans, R. W., Leatham, A. O. and Brooks, R. O. (1985) Powder Melall., 28, 13. Feng, Z., Duszczyk, J. and Leatham, A. G. (1992) Y. Mater. Sci., 27, 4511. Fiedler, H. C., Sawyer, T. E, Kopp, R. W. and Leatham, A. O. (1987) JOM, 39(8), 28. Flemings, M. C. (1974) Solidification Processing. (McGraw-Hill, Inc., New York). Fussell, R. S., Kirchner, H. O. K., Prinz, E B. and Weiss, L. E. (1994) d. Therm. Spray Tech., 3, 148. Grant, P. S. (1995) Prog. Mater. Sci., 39, 497. Grant, P. S. and Cantor, B. (1990) Powder Metall., 33, 144. Grant, P. S., Cantor, B. and Katgerman, L. (1993a) Acta Metall. Mater., 41, 3109. Grant, P. S., Cantor, B. and Katgerman, L. (1993b) Acta Melall. Mater., 41, 3097. Grant, P. S., Kim, W. T. and Cantor, B. (1991) Mater. Sci. Eng., 134A, 1111. Gupta, M., Mohamed, E A. and Lavemia, E. J. (1989) in Proceedings of lhe International Symposium on Advances and Characterization of Ceramic and Metal Malt& Composites, ed. Mostaghaci, H. (Pergamon Press, Oxford, UK), p. 236. Oupta, M., Mohamed, E A. and Lavernia, E. J. (1990) Mater. Manufacturing Proc., 5, 165. Oupta, M., Ibrahim, I. A., Mohamed, E A. and Lavernia, E. J. (1991a) d. Mater. Sci., 26, 6673. Gupta, M., Mohamed, E A. and Lavernia, E. J. (1991b) Internal. d. Rapid Solid., 6, 247. Oupta, M., Mohamed, E A. and Lavernia, E. J. (1992a) Metall. Trans., 23A, 831. Oupta, M., Lane, C. and Lavernia, E. J. (1992b) Scripta Melall. Mater., 26, 825. Gupta, M., Srivatsan, T. S., Mohamed, E A. and Lavernia, E. J. (1993a) Y. Mater. Sci., 28, 2245. Gupta, M., Juarez-Islas, J., Frazier, W. E., Mohamed, E A. and Lavenia, E. J. (1992c) Metall. Trans., 23B, 719. Oupta, M., Bowo, K., Lavernia, E. J. and Earthman, J. C. (1993b) Scripta Melall. Mater., 28, 1053. Outierrez, E. M., Lavernia, E. J., Trapaga, O. M. and Szekely, J. (1988) Inlernat. d. Rapid Solid., 4, 125. Outierrez, E. M., Lavernia, E. J., Trapaga, G. M., Szekely, J. and Grant, N. J. (1989) Metall. Trans., 20A, 71. Hanlon, D. N., Rainforth, W. M. and Sellars, C. M. (1997) Wear, 203/204, 220. Harada, T., Ando, T., O'Handley, R. C. and Grant, N. J. (1991) Mater. Sci. Eng., 133A, 780. Hariprasad, S., Sastry, S. M. L., Jefina, K. L. and Lederich, R. J. (1993) Melall. Trans., 24A, 865. Heimann, R. B. (1996) Plasma Spray Coaling: Principles and Applicalions (VCH Publ. Inc., Weinheim, Germany). Hermann, P. and Morris, D. G. (1994) Metall. Mater. Trans., 25A, 1403. Hermann, P., Biselli, C. and Morris, D. O. (1997) Mater. Sci. Tech., 13, 489. Huron, E. S. (1993) in Spray Forming 2, eds. Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 155. Ibrahim, I. A., Ando, T. and Grant, N. J. (1992) Internal. Y. Rapid Solid., 7, 35. Ibrahim, I. A., Ando, T. and Orant, N. J. (1993) Internal. Y. Rapid Solid., 7, 295. Ibrahim, I. A., Mohamed, E A. and Lavernia, E. J. (1991) d. Mater. Sci., 26, 1137. Ikawa, Y., Itami, T., Kumagai, K., Kawashima, Y., Leatham, A. G., Coombs, J. S. and Brooks, R. G. (1990) ISIY Internal., 30, 756.
Spray-Forming
191
Irani, R. R. and Callis, C. E (1963) Particle Size: Measurement, Interpretation, and Application (John Wiley & Sons, New York). Irmann, R. (1952) Metallurgica, 49, 125. Jones, H. (1978) Aluminum, 54, 274. Juarez-Islas, J. A., Su, S., Perez, R., Campillo, B. and Lavenia, E. J. (1994) Mater. Sci. Eng., 179/180A, 609. Kahl, W. and Leupp, J. (1990) Swiss Mater., 2, 17. Kim, N. J., Park, W. J., Ahn, S. and Elias, L. (1994) in High Performance Metal and Ceramic Composites, ed. Upadhya, K. (TMS, Warrendale, PA), p. 137. Kim, Y. W. (1988) in Dispersion Strengthened Aluminum Alloys, eds. Kim, Y. W. and Griffith, W. M. (TMS, Warrendale, PA), p. 157. Kim, Y. W., Griffith, W. M. and Froes, E H. (1985) JOM, 37(8), 27. Koch, C. C. (1989) Ann. Rev. Mater. Sci., 19, 121. Kurz, W. and Fisher, D. J. (1986) Fundamentals of Solidification. (Trans Tech Publ., Aedermannsdorf, Switzerland). Lavemia, E. J. (1989) Internat. J. Rapid Solid., 5, 47. Lavemia, E. J. and Grant, N. J. (1986) Internat. J. Rapid Solid., 2, 93. Lavemia, E. J. and Grant, N. J. (1988) Mater. Sci. Eng., 98A, 381. Lavemia, E. J. and Wu, Y. (1996) Spray Atomization and Deposition (John Wiley & Sons, Chichester, UK). Lavemia, E. J., McKewan, G. W. and Grant, N. J. (1986) Prog. Powder Metall., 42, 457. Lavemia, E. J., Gomez, E. and Grant, N. J. (1987) Mater. Sci. Eng., 95A, 225. Lavemia, E. J., Srivatsan, T. S. and Mohamed, E A. (1990) J. Mater. Sci., 25, 1137. Lavemia, E. J., Baram, J., and Gutierrez, E. (1991) Mater. Sci. Eng., 132A, 119. Lavemia, E. J., Ayers, J. D. and Srivatsan, T. S. (1992) Internat. Mater. Rev., 37, 1. Lawley, A. (1992) Atomization: The Production of Metal Powders (MPIE Princeton, NJ). Lawley, A. and Apelian, D. (1994) Powder Metall., 37, 123. Lawley, A. and Cantor, B. (1986) in Rapidly Solidified Materials, ed. Lee, E and Carbonara, R. (ASM, Materials Park, OH), p. 15. Lawrynowicz, D. E. and Lavemia, E. J. (1994) Scripta MetaIl. Mater., 31, 1277. Lawrynowicz, D. E., Wolfenstine, J. and Lavemia, E. J. (1997a) Mater. Sci. Eng., 230A, 1. Lawrynowicz, D. E., Li, B. and Lavemia, E. J. (1997b) Metall. Mater. Trans., 28B, 877. Leatham, A. G. (1993) in Spray Forming 2, ed., Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 129. Leatham, A. G. (1996) Advanced Mater. & Proc., 150, 31. Leatham, A. G. and Lawley, A. (1993) Internat. J. Powder Metall., 29, 321. Leatham, A. G., Brooks, R. G. and Yaman, M. (1985) Modem Dev. Powder Metall., 15, 157. Leatham, A. G., Coombs, J. S. and Forrest, J. B. (1995) in Advances in Powder Metallurgy and Particulate Materials, Part 7: Spray Forming, Plasma Processing & SHS, ed. Phillips, M. and Porter, J. (MPIF, Princeton, NJ), p. 31. Lengsfeld, E, Juarez-Islas, J. A., Cassada, W. A. and Lavenia, E. J. (1995) Internat. J. Rapid Solid., 8, 237. Levi, C. G. and Mehrabian, R. (1982) Metall. Trans., 13A, 221. Li, B. and Lavemia, E. J. (1997) Acta Mater., 45, 5015. Li, B., Cai, W. and Lavemia, E. J. (1996a) J. Mater. Synth. Proc., 4, 35. Li, B., Liang, X., Earthman, J. C. and Lavemia, E. J. (1996b) Acta Mater., 44, 2409. Li, B., Nordstrom, N. and Lavemia, E. J. (1997) Mater. Sci. Eng., 237A, 207. Liang, X. and Lavemia, E. J. (1994) Metall. Mater. Trans., 25A, 2341. Liang, X., Earthman, J. C. and Lavemia, E. J. (1992a) Acta Metall. Mater., 40, 3003. Liang, X., Kim, H. K., Earthman, J. C. and Lavemia, E. J. (1992b) Mater. Sci. Eng., 153A, 646. Llorca, J., Martin, A., Ruiz, J. and Elices, M. (1993) Metall. Trans., 24A, 1575. Llorca, J., Ruiz, J., Healy, J. C., Elices, M. and Beevers, C. J. (1994) Mater. Sci. Eng., 185A, 1. Lubanska, H. (1970) JOM, 22(2), 45. Mabuchi, H., Tsuda, H. and Nakayama, Y. (1995) Scripta Metall. Mater., 32, 253. Machler, R., Uggowitzer, E J., Solenthaler, C., Pedrazzoli, R. M. and Speidel, M. O. (1991) Mater. Sci. Tech., 7, 447. Maher, E E, Cantor, B. and Katgerman, L. (1990) in Advanced Aluminum and Magnesium Alloys, eds. Khan, T. and Effenberg, G., p. 659. Majagi, S. I., Ranganathan, K., Lawley, A. and Apelian, D. (1992) in Microstructural Design by Solidification Processing, eds. Lavemia, E. J. and Gungor, M. N. (TMS, Warrendale, PA), p. 139
192
B i n g Li a n d E. J. L a v e r n i a
Marinkovich, J. M., Mohamed, E A., Pickens, J. R. and Lavemia, E. J. (1989) JOM, 41(9), 36. Mathur, E, Apelian, D. and Lawley, A. (1989) Acta Metall., 37, 429. McLelland, A. R. A., Atkinson, H. V., Kapranos, E and Kirkwood, D. H. (1991) Mater. Lett., 11, 26. Megusar, J., Lavernia, E. J., Domalavage, P., Harling, D. K. and Grant, N. J. (1984) J. Nucl. Mater., 122/123A, 789. Morris, D. G., Morris, M. A., Gunter, S., Leboeuf, M. and Hollrigl, G. (1992) Scripta Metall. Mater., 27, 1645. Namba, Y., Saito, M., Matsushita, T. and Takigawa, H. (1993) in Spray Forming 2, ed. Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 221. Noguchi, A., Ezawa, I., Kaneko, J. and Sugamata, M. (1995) J. Japan Inst. Light Metals, 45, 64. Ogata, K., Lavemia, E. J., Rai, G. and Grant, N. J. (1986) lnternat. J. Rapid Solid., 2, 21. Palmer, I. G., Chellman, D. J. and White, J. (1993) in Spray Forming 2, ed. Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 385. Parker, D. W. and Kutner, G. L. (1991) Advanced Mater. & Proc., 139, 68. Perez, J. E and Morris, D. G. (1994) Scripta Metall. Mater., 31, 231. Perez, R. J., Zhang, J. and Lavemia, E. J. (1992) Scripta Metall. Mater., 27, 1111. Perez, R. J., Zhang, J., Gungor, M. N. and Lavernia, E. J. (1993) Metall. Trans., 24A, 701. Perlinde, G. and Lutjering, G. (1982) Metall. Trans., 13A, 1283. Picraux, S. T. (1984) Ann. Rev. Mater. Sci., 14, 335. Ranz, W. E. and Marshall, W. R. (1952) Chem. Eng. Prog., 48, 141. Rawers, J. C. and Wrzesinski, W. R. (1992) J. Mater. Sci., 27, 2877. Reed-Hill, R. E. (1973) Physical Metallurgy Principles (2nd edition, Von Nostrand, Princeton, NJ). Report, "Spray Deposited High Performance Materials: N202", Alusuisse-Lonza services Ltd. (contact person: W. Kahl), June, 1990a. Report, "Spray Deposited High Performance Materials: N403", Alusuisse-Lonza services Ltd. (contact person: W. Kahl), June, 1990b. Report, "Spray Deposited High Performance Materials: N707", Alusuisse-Lonza services Ltd. (contact person: W. Kahl), June, 1990c. Reppich, B. (1993) in Plastic Deformation and Fracture of Materials, ed. Mughrabi, H., Vol 6 of Materials Science and Technology (VCH Publ. Inc., Weinheim, Germany), p. 311. Rickinson, B. A., Kirk, E A. and Davies, D. R. G. (1981) Powder Metall., 24, 1. Ruhr, M. and Baram, J. (1991) Metall. Trans., 22A, 2503. Ruhr, M., Lavemia, E. J. and Baram, J. (1990) Metall. Trans., 21A, 1785. Sampath, S., Tiwari, R., Gudmundsson, B. and Herman, H. (1991) Scripta Metall. Mater., 25, 1425. Singer, A. R. E. (1970) Met. and Mater., 4, 246. Singer, A. R. E. (1972) J. Inst. Met., 100, 185. Singer, A. R. E. (1982) Powder Metall., 25, 195. Singer, A. R. E. and Evans, R. W. (1983) Met. Technol., 10, 6. Singer, A. R. E. and Ozbek, S. (1985) Powder Metall., 28, 72. Singh, R. E, Lawley, A., Friedman, S. and Nurty, Y. V. (1991) Mater. Sci. Eng., 145A, 243. Song, G., Shen, J., Jiang Z. and Li, Q. (1994) Mater. Sci. Eng., 179/180A, 249. Srivatsan, T. S. and Lavernia, E. J. (1994) Composite Eng., 4, 459. Steen, W. M. (1985) Met. and Mater., 1, 730. Tanigawa, S., Hara, Y., Lavemia, E. J., Chin, T. S., O'Handley, R. C. and Grant, N. J. (1986) IEEE Trans. Mag., 22, 746. Turnbull, D. and Cech, R. E. (1950) J. Appl. Phys., 21, 804. Ucok, I., Ando, T. and Grant, N. J. (1991) Internat. J. Powder Metall., 27, 237. Underhill, R. E, Grant, E S., Bryant, D. J. and Cantor, B. (1995) J. Mater. Synth. Process., 3, 171. Underhill, R. U., Grant, E S. and Cantor, B. (1993) Mater. Design, 14, 45. Upadhyaya, A., Mishra, N. S. and Ojha, S. N. (1997) J. Mater. Sci., 32, 3227. Veistinen, M. K., Hara, Y., Lavemia, E. J., O'Handley, R. C. and Grant, N. J. (1987) in High Performance Permanent Magnet Materials Symposium, eds. Sankar, S. G., Herbst, J. F. and Koon, N. C. (MRS, Pittsburgh, PA), p. 93. Vetter, R., Zhuang, L. Z., Majewska-Glabus, I. and Duszezyk, J. (1990) Scripta Metall. Mater., 24, 2025. White, J., Mingard, K., Hughes, I. R. and Palmer, I. G. (1993) in Spray Forming 2, ed. Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 355. White, J., Mingard, K., Hughes, I. R. and Palmer, I. G. (1994) Powder Metall., 37, 129.
Spray-Forming
193
White, J., Willis, T. C., Hughes, I. R. and Jordan, R. M. (1988) in Dispersion Strengthened Aluminium Alloys, eds., Kim, Y. W. and Griffith, W. M. (TMS, Warrendale, PA), p. 693. White, J., Palmer, I. G., Hughes, I. R. and Court, S. A. (1989) in Aluminum-Lithium Alloys, eds. Sanders, Jr., T. H. and Starter, Jr., E. A. (MCE Publ., Birmingham, UK), p. 1635. Willis, J. (1988) Met. and Mater., 4, 485. Wu, M., Srivatsan, T. S., Pickens, J. R. and Lavernia, E. J. (1992) Scripta Metall. Mater., 27, 761. Wu, Y. and Lavemia, E. J. (1991) JOM, 43(8), 16. Wu, Y. and Lavernia, E. J. (1992) Metall. Trans., 23A, 2923. Wu, Y., Cassada, W. A. and Lavernia, E. J. (1993) in International Conference on Advanced Synthesis of Engineered Structural Materials, eds. Moore, J. J., Lavernia, E. J. and Froes, E H. (ASM, Materials Park, OH), p. 245. Wu, Y., Zhang, J. and Lavernia, E. J. (1994) Metall. Mater Trans., 25B, 135. Wu, Y., Cassada, W. A. and Lavemia, E. J. (1995) Metall. Mater Trans., 26A, 1235. Wu, Y., Del Castillo, L. and Lavemia, E. J. (1997) Metall. Mater Trans., 28A, 1059. Yule, A. J. and Dunkley, J. J. (1994) Atomization of Melts for Powder Production and Spray Deposition (Clarendon Press, Oxford, UK). Zeng, X., Liu, H., Chu, M. G. and Lavernia, E. J. (1992) Metall. Trans., 23A, 3394. Zeng, X., Nutt, S. R. and Lavernia, E. J. (1995) Metall. Mater Trans., 26A, 817. Zhang, J., Perez, R. J., Gupta, M. and Lavernia, E. J. (1993a) Scripta Metall. Mater., 28, 91. Zhang, J., Gungor, M. N. and Lavernia, E.J. (1993b) J. Mater Sci., 28, 1515. Zhang, J., Perez, R. J. and Lavernia, E. J. (1994) Acta Metall. Mater., 42, 395. Zhang, Y. C., Wu, Y., Zhao, Y. Y., Liu, Z. E and Chen, G. Y. (1989) inAluminium-Lithium Alloys, eds. Sanders, Jr., T. H. and Starter, Jr., E. A. (MCE Publ., Birmingham, UK), p. 95. Zhao, Y. Y., Grant, P. S. and Cantor, B. (1993a) J. Physique, 3, 1685. Zhao, Y. Y., Grant, P. S. and Cantor, B. (1993b) J. Microscopy, 169, 263. Zheng, J. H. and Kahl, W. (1990) in Advanced Aluminum and Magnesium Alloys, eds. Khan, T. and Effenberg, G., p. 317. Zhou, J., Duszczyk, J. and Korevaar, B. M. (1991) J. Mater. Sci., 26, 5275. Zhuang, L. Z., Majewska-Glabus, I., Vetter, R. and Duszczyk, J. (1990) Scripta Metall. Mater., 24, 2025.
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Chapter 8 Ion-Mixing 8.1. 8.2. 8.3.
Introduction Brief Description of Underlying Physics in Ion Mixing Thermodynamics of Alloy Phase Formation 8.3.1 Miedema's Theory and Alonso's Method 8.3.2 Interfacial Free Energy in the Multilayers 8.4. Experimentation of Ion-Mixing 8.4.1 Sample Design 8.4.2 Ion-Mixing Parameters 8.4.3 Characterization Methods 8.5. Amorphous Phase Formation 8.5.1 Glass-Forming Ability of Systems with a Negative Heat of Formation 8.5.2 Glass-Forming Ability of Systems with a Positive Heat of Formation 8.5.2.1 Amorphous Alloys Formed Within Restricted Compositions 8.5.2.2 Amorphous Alloys Formed in a Broad Composition Range 8.5.2.3 Nominal and Intrinsic Glass-Forming Ability 8.6. Formation of Metastable Crystalline Alloys 8.6.1 Structural Classification of the Metastable Crystalline Phases 8.6.1.1 Solid Solutions 8.6.1.2 h.c.p.-I and f.c.c.-I Phases 8.6.1.3 h.c.p.-II and f.c.c.-II Phases Based on b.c.c. Metals 8.6.2 Free Energy calculation of the MX phases 8.7. Interface-Generated Solid-State Vitrification in Systems with a Positive Heat of Formation 8.8. Concluding Remarks Acknowledgments References
197 199 200 200 201 203 203 205 205 205 206 207 207 207 211 213 213 213 213 215 216 218 219 220 220
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Chapter 8 Ion-Mixing B. X. LIU
8.1. I N T R O D U C T I O N
The earliest use of an energetic ion beam generated from low-energy accelerators (several 100-keV or higher) to modify material properties was initiated in early 1960s, i.e. the doping of semiconductors by implanting the desired ion species into silicon wafers (Carter and Grant 1976). Since then, ion implantation has significantly developed and has become a supporting technique in the semiconductor and VLSI technology (Dearnaley et al. 1973). Since early 1970s, ion-implantation has also been employed to deal with metallic materials, e.g., for improving the wear and corrosion resistance, catalytic properties, superconductivity, etc. (Mayer and Lau 1988). As a versatile metal ion source has not been available for some time, it has not been possible to implant various metal species to achieve a high ion concentration frequently required for modifying the surface properties of the metallic materials. An alternative scheme, namely ion-mixing of multiple metal layers, was therefore introduced and a new research area of synthesizing metastable alloys by ion-mixing has boomed since the early 1980s (Liu et al. 1983). Figure 8.1 shows schematically how ion-mixing is realized in multilayered films consisting of metal layers A and B. The A-B multilayered films are prepared by depositing alternately the metals A and B on a substrate. The as-deposited films are then irradiated by a scanning ion beam, e.g., 300 keV xenon ions, to a certain dose to achieve a uniform mixing between metal layers A and B, resulting in forming an Al-xBx alloy. The substrate can be SiO2 or NaC1 acting as a supporting substance or it can be metal A as a matrix for forming a surface Al-xBx alloy coating on metal A. It is known that the first amorphous alloy, also called metallic glass, was obtained by liquid melt quenching (LMQ) in 1960 (Klement et al. 1960). Because amorphous alloys have a non-crystalline structure (with no vulnerable grain boundaries), they are expected to possess unique properties. During the past 30 years, various techniques have been developed to synthesize metallic glasses. From a physical point of view, high cooling rates of 106 K/s or higher are needed to avoid crystallization against amorphization (Davies 1983, Hafner 1981, Giessen and Whang 1980) and cooling rates of up to 107 K/s have been achieved so far. Such developments lead to expand the number and types of metallic glasses, including a variety of metal-metalloid and metal-metal amorphous alloys (Takayama 1976). Laser-quenching has also been applied for producing amorphous alloys (von Allmen 1983, Lin and Spaepen 1984), as many intense laser sources have 197
198
B. X. Liu
amorph~s SiOz
... ..
: NaCI 3,.keVXeions
:~~7
[
cn'statline Figure 8.1. Ion-mixingof multilayered films consisting of metal layers A and B.
become available and they can achieve a cooling rate of 1012 K]s in the material's surface layer (See the chapter on "Laser Processing" by K.E Kobayashi in this book). In comparison, effective cooling rates can reach as high as 1014 K]s (Thompson 1969) in ion-mixing, and therefore this technique can also produce a variety of amorphous alloys. In fact, the cooling rates achieved during ion-mixing are probably the highest among the currently available glass producing techniques. Up-to-now, numerous metastable alloys with either an amorphous or crystalline structure have been obtained by ion-mixing in some 100 binary metal systems (Liu and Jin 1997), and in many alloy systems that are not obtainable by other non-equilibrium processing techniques. The glass-forming ability (GFA) of a system can be considerably improved by ion-mixing because of its high cooling rate. Furthermore, ion-mixing offers a unique opportunity for studying the amorphization mechanism, since the ion dose can be varied at small intervals and so the process of ion-mixing can be studied step-by-step. Based on experimental data, several empirical models have been proposed to predict the GFA and the glass-forming composition range (abbreviated hereafter as GFR) of the binary metal systems. Liu et al. (1983) proposed the first rule, namely the structural difference rule, which predicted a sufficient condition and demonstrated the predominant role of the crystalline structures of the component metals in forming metallic glasses by ionmixing. Further studies led Liu to formulate a two-parameter model. The first parameter is the maximum possible amorphization range (MPAR), which was defined as the total width of the two-phase region, equaling the whole composition range (100%) minus the maximum solid solubilities, observed from the corresponding equilibrium phase diagram (Liu et al. 1987). The second parameter is the heat of formation A H f calculated by Miedema's model (Miedema 1976). Based on a quantitative combination of the parameters, the binary metal systems are classified into readily, possibly, and hardly glass-forming ones (Liu et al. 1987). Alonso and Simozar (1983) constructed a two-dimensional map with
199
Ion-Mixing
the ratio of atomic radii, rA/rB, and the A H f and by summarizing the published data, they separated the binary metal systems into glass-forming and non-glass-forming ones and predicted that no amorphous alloy could be formed in a system with a A H f exceeding + 10 kJ/mol. Ossi (1990) proposed an atomistic model based on collision cascades and the related bombardment-induced surface compositional changes. Though the prediction by the two-parameter model matched best with the experimental data, lack of a relevant thermodynamic interpretation was the main drawback until the early 1990s. It was also noted that metallic glasses were unexpectedly obtained in some alloy systems with a positive A Hu, which have long been recognized as non-glass-forming ones (Liu 1987, Jin et al. 1995, Zhang et al. 1995). The thermodynamic stability of the different alloy phases was evaluated by Miedema, who developed a semi-quantitative model in the mid 1970s to calculate the free energy curves of the alloy phases versus alloy compositions by using the available characteristic parameters of the constituent metals (de Boer et al. 1989), rather than the parameters of the alloy phases, as required in CALPHAD calculation. This chapter attempts to present a brief review of the up-to-date progress on the synthesis of metastable alloys by ion-mixing of multilayered films. It is divided into the following sections: a brief description of the underlying physics in ion-mixing, thermodynamic calculations based on Miedema's model, experimentation and sample design, formation of metallic glasses in systems with a positive or a negative heat of formation, metastable crystalline alloys synthesized by ion mixing, and applicability of the interface concept to solid-state amorphization.
8.2. B R I E F D E S C R I P T I O N O F U N D E R L Y I N G P H Y S I C S IN I O N M I X I N G
This section briefly describes the underlying physics in ion-mixing and the experimental details of ion-mixing will be discussed in Section 8.4. The ion-mixing process is commonly divided into two steps, viz., an atomic collision cascade and subsequent relaxation. After relaxation, one can consider having a delayed step and it has no significant effect on alloy phase formation. When an ion with an energy E1 impacts on the material surface and penetrates inside, it collides with the atoms in the material. If an atom in the material receives an energy E2 exceeding a threshold energy Ed, the atom will leave its own lattice site and a vacancy is created. This displaced atom is referred to as knock-out atom and it will, if E2 is greater than Ed, trigger secondary collisions, then tertiary ones, and so on, and this is named an atomic collision cascade. To describe atomic displacements and defect creation caused by ion irradiation, a parameter "dpa", (abbreviation for displacement p e r atom) is defined, which can be calculated through atomic collision theory, as
do-
E2 (max)
d p a -- Nd = ~ . Crc
No
rJed
v(E2)
9dE2 ~
(8 1) "
whereNd and No are the numbers of displaced and total atoms per volume, respectively; r is the ion flux, ere is the collision cross-section; Ed and E2 are the displacement threshold energy and the energy received by the knock-out atom, respectively; v(E2) is the number
200
B. X. Liu
of displaced atoms created by the knock-out atom with an energy of E2, and dot/dE2 is the differential collision cross-section at E2. The atomic collision cascades trigger intermixing of the metal layers A and B in the multilayered films. As the energy of the ions is on the order of several 100 keV and the binding energy of the solids is typically 5-10 eV, atomic collision is a highly dynamic process of far-from equilibrium. By receiving an adequate irradiation dose, the layered structure of the multilayered films is smeared out and metals A and B mix into a uniform mixture, which is in a highly energetic state with a disordered atomic configuration. At the termination of the cascade, the highly energetic mixture should somehow relax towards equilibrium, and at this stage, the equilibrium thermodynamics comes into play in governing the direction of relaxation. Whether the mixture can reach the equilibrium state or not, however, depends on the temperature and time conditions available during this period. During the atomic collision cascades, all the atoms are involved in violent motion, it is hard to imagine the formation of any kind of alloy phase. Consequently, it is commonly considered that the structure of the alloy phase is to be formed during the relaxation period. As the relaxation period is very short (10-1~ -9 seconds, as we will discuss soon) and the mixture cannot straightforwardly go to the equilibrium state in most cases, it frequently resides in some intermediate states, corresponding to the metastable states of either an amorphous or a simple-structured crystalline phase. From an energy point of view, a thermal spike concept was introduced to discuss the ion-mixing process. Suppose a target is irradiated by ions of 300 keV energy and a current density of 1-4/zA/cm 2 for a certain time period, a huge amount of energy would be deposited into the material surface within a very thin layer. Estimation showed that in a local region, e.g., using a cylinder geometry, with a diameter of 5 nm and a length of 100 nm, the temperature could be as high as 103-104 K, while the surrounding material remains close to the ambient temperature of 300 K. According to thermal conduction calculations, the temperature drops to 300 K within approximately 10-1~ -9 second, which is the duration of relaxation. Apparently, during relaxation the effective cooling rate can be as high as 1014 K/s (Thompson 1969).
8.3. T H E R M O D Y N A M I C S
OF A L L O Y P H A S E F O R M A T I O N
As described above, the process of ion-mixing, as a whole, takes place under conditions that are far from equilibrium. However, during the relaxation period, the highly energetic state relaxes towards equilibrium. At this stage, equilibrium thermodynamics plays a role in governing the possible evolution path, thus determining the nature (amorphous or crystalline) of the alloy phase formed. It is therefore necessary to calculate the free energy of the competing phases in the alloy system to understand and predict alloy phase formation upon ion-mixing.
8.3.1 Miedema's theory and Alonso's method In this section, the discussion is focused on the amorphous phase formation and the thermodynamics for metastable crystalline phase formation will be discussed later.
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Ion-Mixing
Generally, the Gibbs free energy A G of a phase is calculated by A G = A H -- T A S, where A H and AS are the enthalpy and entropy of formation, respectively. The entropy usually plays only a secondary role in solid phase formation, and its expression can be taken, as a first approximation, as that of an ideal solid solution, i.e., AS = - - R [ x A ln(xa) + XB ln(XB)], where R is the gas constant, and XA and XB are the atomic concentrations of metals A and B, respectively. For a substitutional solid solution of transition metals, the enthalpy of formation, according to Miedema's model and Alonso's method (Alonso et al. 1990, Niessen and Miedema 1983, Lopez and Alonso 1985), can be considered as a sum of three contributions A H = A H c + A H e + A H s,
(8.2)
where A H c, A H e and A H s correspond to the chemical, elastic, and structural contributions, respectively, and they can be calculated according to the well documented theory and methods. For an amorphous phase, the enthalpy of formation due to both elastic and structural terms is absent. Following the suggestion of Niessen et al. (1988), the enthalpy of the amorphous phase is then given by
(8.3)
AHamorphous-- A H c @OI(XATm,A @ XBTm,B),
where ot is an empirical constant equal to 3.5 J/(mol 9K) and Tm,i is the melting point of the component i. 8.3.2 I n t e r f a c i a l f r e e energy in the multilayers
Since ion-mixing (IM) begins with multilayered films, which definitely consist of a number of interfaces, the effect of interfacial free energy on the alloying behavior should be taken into account. To include the effect of interfaces, the Gibbs free energy of the initial state of the A-B multilayered films was calculated by adding the interfacial free energy to the ground state of a mixture of A and B in bulk form. As pointed out by Turnbull (1955), the interfacial free energy between two solid phases can be split into geometrical and chemical contributions and can be written as ysAB m
.geo .chem r s s + Yss 9
(8.4) geo For two solids with quite different lattice structures, the geometrical term Yss stands for the energy of a large-angle grain boundary. It is estimated that the grain boundary enthalpy is about 30 percent of the surface energy at 0 K (Murr 1975) and the entropy contribution is neglected. Thus .
.chem , y~sAsB - - 0 . 1 5 x (ys,o + YBB,0) + YSS
(8.5)
where yjS,0, ys,0 are the surface energies of metals A and B, respectively, and Yss'chem is given by Miedema (1978)
B. X. Liu
202 ----d
---'B
--dA'-~ ~ d B - ~ d A
B
A
I I I I I I
iI I I I I I I
A
Figure 8.2. Schematicconfiguration of A-B multilayered films.
ychem
SS
i12/3
-- AnOinB/(CO, a ),
(8.6)
where Co is a constant being 4.5 • 108, A H ~ in B is the heat of solution of metal A in metal B at infinite dilution (Miedema 1978) and VA is the atomic volume of metal A. According to Gerkema and Miedema (1983), the interfacial free energy of the multilayers, AGmul was calculated as follows"
A Gmul - S f • ysAsB,
(8.7)
where Sf is the surface area occupied by one mole of interfacial atoms. Figure 8.2 shows schematically the configuration of the A-B multilayers. Let us assume that NA and N8 are the numbers of layers of metals A and B, dA,i and dB,i are the thicknesses of the ith layers of metals A and B, respectively; nA and n8 are numbers of interfaces (A on B) and (B on A); AdA,i and AdB,i are the thicknesses of the ith interfaces (A on B) and (B on A), respectively. Obviously, the relationship of n A+ n 8 + 1 = NA +N8 always holds for these four parameters. The interfacial free energy of the multilayers can then be calculated as
(8.8)
AGmul -- OtASf AYBA + OIBSfBYAB,
where SfA and SfB are the surface areas occupied by one mole of A and B atoms, respectively, lYa and OrB are the ratios of interfacial atoms A and B to the total number of atoms in the multilayered films, respectively, and are given by (Zhang and Liu 1995)
S nA AdA,i OlA -- XA Z NA dA,i
Z j B mdB,j ,
or8 - x 8
~NB NB,j
,
(8.9)
Z ; A AriA, i / ~-~NA dA,i and )--~.8 A d s , j / ~--~8 ds,j are the ratios of interfacial A atoms to the total number of A atoms and the ratio of interfacial B atoms to the total number of B atoms, respectively. These terms, multiplied by XA and x s, respectively, give the fractions of A and B interfacial atoms versus the total number of (A + B) atoms where
Ion-Mixing
203
in the multilayers. According to experimental studies on thin films (Clemens 1986) and for simplicity of calculation, the thickness of each interface is assumed to have the same value of 0.5 nm. It should be pointed out that in experimental studies one can always add a few extra interfaces in a sample design to ensure a higher energetic state of the multilayers than that of the phase of interest, so that the assumption of the effect of this thickness can be ignored. Figure 8.3 shows two free energy diagrams calculated by the above methods for two representative systems with negative (Ni-Nb) and positive (Y-Nb) AHf to show a significant difference due to a different sign of the heat of formation. One sees that the free energy curve of the amorphous phase for the Ni-Nb system is concave, i.e., it is always lower than that of an equilibrium mixture of two crystalline metals, suggesting a favored condition for forming glass. The free energy curve of the amorphous phase in the Y-Nb system is convex and is always higher than that of the equilibrium state without considering the interfacial free energy. It is noted that taking into account the interfacial free energy, the energetic sequence of the respective states in the Y-Nb system is changed significantly, while the effect of interface is negligible for the Ni-Nb system. The details of this will be further discussed together with the experimental results later.
8.4. E X P E R I M E N T A T I O N
OF I O N - M I X I N G
8.4.1 Sample design In designing the multilayered samples, the total thickness of the multilayers should match with the range of the irradiating ions. Inert gas ions, such as Xe, Kr, Ne, or Ar, are frequently chosen to be the irradiating ions. The advantages of employing inert gas ions are that the chemical impurity effect is negligible, if the ion dose is below 1016 ion/cm 2, and that the gas ion sources are frequently available easily for the common implanters. As heavier ions induce more efficient mixing than the lighter ones when the other conditions are kept the same, Xe + ions are therefore mostly used. The typical energy of the irradiating ions is several 100 keV. It was found experimentally that if the total thickness of the multilayered film was designed to be the projected range Rp plus the projected range straggling ARp of the irradiating ions, uniform mixing between the layers could be achieved well at an ion dose range of 1014-1016 ion/cm 2. The Rp and ARp for elemental metals can be calculated easily by the LSS theory (Dearnaley et al. 1973). For an alloy target of AxB 100-x, where x refers to the atomic percentage of metal A, the projected range of the alloy can be calculated by the following equation (Townsend et al. 1976):
Rp (Ax B 100-x ) --
Rp (A) Rp (B) yRp(B) + ( 1 0 0 - x)Rp(A)'
(8.10)
where y refers to the weight percentage of metal B in the alloy and can be obtained from x easily. The calculation of the projected range straggling ARp for the alloys is rather complicated. However, considering the errors involved in the calculation itself, as well as in controlling the thickness during deposition, the above equation is also good enough to estimate ARp of the alloy, by substituting ARp and Rp. The total thickness is divided into the thicknesses of metals A and B, according to the desired atomic ratio of the alloy being
204
B. X. Liu
Nb
Y ,
,
(a)
.
,
Muitilayers(17 interfaces)
+60
Amorphous +40
+20
0
&
+20
&
I
9
Multilayers(7 interfaces) (17 interfaces) --"-'--'"'"'""-7~.~ --
-20
----7,
Amorphous
(b) Ni
I
20
9
I
1
40 60 80 Nb concentration (at. %)
J
Nb
Figure 8.3. Calculated free energy diagrams of the Y-Nb (a) and Ni-Nb (b) systems showing the different effects of interfacial free energy on the energetic sequence in the positive and negative heat of formation systems.
studied. Typically, the total thickness of the films of transition metal alloys is about 60100 nm under 300 keV xenon ion-mixing. Another parameter is the fraction of interfacial atoms, which determines the number of metal layers. Each individual layer is always designed to be 5-8 nm or even thinner. The multilayered films are prepared in a vacuum system to deposit alternately the metals A and B onto an inert substrate supporting the film or on a metal matrix. The supporting materials can be SiO2, A1203 or NaC1 single crystals. The system is often a multi-crucible electron gun evaporation system with a vacuum level better than 10 -5 Pa.
Ion-Mixing
205
8.4.2 lon-mixing parameters It was found experimentally that the ion dose required to homogenize the discrete layered films is in the range of 1014 to 1016 ions/cm 2 and varies from system to system. The ionbeam current density should be controlled generally around 1/~A/cm 2, or at most up to 4/zA/cm 2 to avoid overheating of the multilayers during ion irradiation. In operating the implanter, the ion dose, which is measured by the electrical charge, is easy to adjust and control. Ion-mixing thus offers a unique method for investigating the phase transformation process step-by-step. Another important parameter in ion-mixing is the temperature at which the experiment is conducted. As ion-mixing often forms metastable alloy phases, ion-irradiation is generally performed at low temperatures, e.g., at room temperature or down to liquid nitrogen temperature. Nevertheless, the temperature mentioned for the experiments is often the nominal temperature measured at the sample holder instead of the effective temperature of the irradiated films. The effective temperature is an important parameter in the process induced by ion irradiation, and it should be considered when discussing the alloy phase formation mechanism. The stable temperature of the expecting alloy should be considered, i.e., if one expects to form an amorphous alloy with a recrystallization temperature Ta-c (from amorphous-to-crystalline), the effective temperature should of course be kept at a temperature far below Ta-c.
8.4.3 Characterization methods 1-2 MeV He + ion Rutherford backscattering spectrometry (RBS) (Chu et al. 1978) is frequently used to analyze the amount of mixing in the multilayered films. Figure 8.4 shows a typical example for Fe-Mo multilayered films before and after ion-mixing and one can clearly see the information given by the RBS spectra. Besides, XPS, SIMS and Auger spectroscopy are employed to analyze the composition depth profile of the ion mixed films. The structure of the phase is identified mainly by X-ray diffraction and TEM. Fourpoint probe resistivity measurement is a supplemental method to give evidence that the amorphous phase is formed by a crystal-to-amorphous transformation, because the resistivity of the amorphous alloys is often considerably greater than that of their crystalline counterparts.
8.5. A M O R P H O U S
PHASE FORMATION
The maximum possible amorphization range (MPAR) is a decisive parameter in determining the GFA or GFR. When a system has an MPAR exceeding 20%, the amorphous phase can probably be formed in the system. If the MPAR is less than 20%, the related two-phase region is too narrow and can be further decreased by the extension of the solid solubilities from two metal ends under ion-mixing. In other words, it is difficult to form an amorphous alloy (Liu 1987). For an extreme value of MPAR = 0, the phase diagram of the system is of continuous solid solution type with no two-phase region, and GFA and GFR are zero. The following discussion will show that the empirical prediction agrees
206
B. X. Liu
~.~
Mo
2.0MeV
SiO2 ~ ~~70.
He
Fe
+
Mo
.,4 ~a 4~
Ix 1016Xe,./ c m
2
\
as-deposited'
|
. ~ _
500
i_ =.........
600
Channel. Number Figure 8.4. Backscattering spectra of Fe-Mo multilayered films before (solid line) and after room temperature ion-mixing to a dose of 1 • 1016 Xe+/cm 2.
reasonably well with the experimental results for the systems featuring either negative or positive A H f 8.5.1 Glass-forming ability of systems with a negative heat of formation A large body of data confirmed that an amorphous alloy is most likely to form by IM in the two-phase region (Liu et al. 1983), which can be attributed to the competition of crystallization between two crystalline phases, resulting in a frustration of nucleation and growth of the crystalline phase. It is also easy to form amorphous alloys at or in the vicinity of intermetallic compound compositions. This is even more so when the intermetallics have complex crystal structures and therefore cannot form due to a very restricted kinetic condition available in IM (Liu and Jin 1997). Consider the Ni-Nb system (A H i = - 4 2 kJ/mol) as a representative example, whose calculated free energy diagram is shown in Figure 8.3. One sees clearly that the energetic sequence of the phases involved is thermodynamically favorable for the formation of a glassy phase. As the formation of intermetallic compounds is usually hindered kinetically, they need not be considered when discussing the GFA. The GFR was therefore regarded as the composition range in which the free energy of the amorphous phase is lower than
Ion-Mixing
207
that of the solid solutions. The points of intersection between the free energy curve of the amorphous phase with those of the two solid solutions are around 20 and 85 at. % Ni. The GFR in Ni-Nb is therefore from 20 to 85 at.% Ni, which is in excellent agreement with the IM results, i.e. 20-85 at.% Ni, observed by Zhang et al. (1993), as well as very close to the defined MPAR of the system.
8.5.2 Glass-forming ability of systems with a positive heat of formation For the Y-Nb system with a positive A H i, shown in Figure 8.3, the free energy curve of the amorphous phase is usually higher than that of an equilibrium mixture of crystalline metals A and B in bulk form, leading to a situation unfavorable for amorphization. However, as the interfaces in multilayered films are generally regarded as transient layers of mixed A and B atoms in a highly disordered configuration, they contribute extra free energy to the films. With an increasing fraction of the interfacial atoms, it is possible to have the free energy curve of the multilayers intersect with that of the amorphous phase in some specific composition ranges, making it possible to form a glass. These composition ranges depend on the designed fraction of the interfacial atoms of the films (Zhang et al. 1995). Some examples of experimental results are given in the following paragraphs.
8.5.2.1 Amorphous alloys formed within restricted compositions Figure 8.5 shows the calculated free energy diagram of the Ag-Mo system with a A Hf of +56 kJ/mol, the second largest A H f for transition metal binary systems. One sees that the free energy curve of the multilayered films containing 12 layers (or 11 interfaces) intersects with that of the amorphous phase, dividing the diagram into three regions along the composition axis. Obviously, in regions I and III, the amorphous phase is possible to be formed, while in region II, amorphization is hardly to be achieved. Ion-mixing by 200 keV xenon ions at liquid nitrogen temperature did result in the formation of an amorphous alloy at near 80 at.% Mo, and the formation of another amorphous phase co-existed with a new Mo-based f.c.c, phase around 40 at.% Mo (Jin et al. 1995). Such compositional discontinuity in glass formation has also been observed in other systems with a positive A H f (Jin and Liu 1994).
8.5.2.2 Amorphous alloys formed in a broad composition range With a further increase in the fraction of interfacial atoms versus the total number of atoms in the multilayered films, the two separated composition ranges favoring amorphization may become broader and meet each other to become a continuous one. The calculated free energy diagram of the Y-Mo system with A H f = +35 kJ/mol, shown in Figure 8.6, includes three free energy curves for the multilayered films with 9, 11 and 19 layers, corresponding to the fractions of interfacial atoms of 8.1%, 10.0% and 18.1%, respectively. Table 8.1 lists the predicted composition ranges favoring amorphization for these three cases. The Y-Mo multilayered films were prepared accordingly for ion-mixing investigation and the experimental results are in excellent agreement with the prediction (Zhang et al. 1995). These results suggest that through appropriately designing the multilayers, it is possible to synthesize metallic glasses in systems with positive A Hf even with a composition near the equiatomic stoichiometry, where the amorphous phase frequently has the
208
B. X. Liu
_~-
O
6O
>'
40
"
i
~
I
'
I
'
I
E
30 ii
~
2o
~
10
L_
,
f.c
0
12 L a y e r s
)
(.9
I
~
v
K,.,
'
-
rphous
1
l
Mo 10
1
,
1
30
50
9
1
70
j.
~'~
I
90
Ag
Ag C o n c e n t r a t i o n Figure 8.5. Calculated free energy diagram of the A g - M o system.
Table 8.1 Predicted and experimental results of amorphous phase formation in the Y-Mo multilayered films upon 200 keV xenon ion mixing at room temperature, n" number of interfaces in the multilayers; o~" fraction of interfacial atoms in the multilayers. Composition
Y80 Mo20
Y60 Mo40
Y50 Mo50
Y40 Mo60
Y20 Mo80 Yes
n = 18
Predicted
Yes
Yes
Yes
Yes
(a = 18.1%)
Ion mixing
Yes
Yes
Yes
Yes
*
n = 10
Predicted
Yes
Yes
No
No
Yes
(or = 10%)
Ion mixing
Yes
Yes
No
No
Yes
n =8
Predicted
Yes
No
No
No
Yes
ot = 8.1%)
Ion mixing
Yes
No
No
No
Yes
* Experiment not conducted.
highest free energy and is most difficult to form. It is worth mentioning that the IM results listed in Table 8.1 include both "yes" and "no" and that they strongly depend on the correlation between interfacial fraction and alloy composition. These results also support the argument that IM induced amorphization can really be attributed to the interfacial free energy, but not to the irradiation energy. The irradiation energy has exactly the same effect on all the multilayered films and does not depend on alloy composition. In the past 15 years, nearly 100 binary metal-metal systems, according to published results, have been studied by ion-mixing of multilayers. Table 8.2 lists these systems in alphabetical order of the chemical symbols of the constituent metals for convenience of searching.
Ion-Mixing
209
Table 8.2 Binary metal-metal systems that have so far been studied by ion mixing AHf
Any
<0
System
Ref.
Remarks
A1-Au A1-Co A1-Fe A1-Mn A1-Mo A1-Nb A1-Ni A1-Pd A1-Pt A1-Ta A1-Ti A1-Y A1-Yb Au-Ta Au-Ti Au-V Co--Gd Co-Hf Co-Mo Co-Nb Co--Ta Co-Tb Co-Ti Co-Y Co-Yb Co-Zr Cu-Y Cu-Zr Er-Ni Fe-Gd Fe-Mo Fe-Nb Fe-Ta Fe-Tb Fe-Ti Fe-Y Fe-Zr Hf-Ni Mo-Ni Mo-Ru Mo-Zr Nb-Ni Ni-Ta Ni-Ti Ni-Y Ni-Yb Ni-Zr Ru-Ti Ru-Zr
Schmid and Ziemann (1985) Hung and Mayer (1985) Rauschenbach and Hohmuth (1987) Knapp and Follstaedt (1987) Meissner et aL (1987) Liu et al. (1983) Hung et al. (1983) Hung et al. (1983) Hung et al. (1983) Meissner et al. (1987) Cheng (1986) Liu et al. (1991) Ding et al. (1992) Pan et aL (1996) Liu et al. (1982) Tsaur et aL (1991) Yan et aL (1986) Anderson et al. (1990) Liu et al. (1983) Zhang and Liu (1994) Zhang and Liu (1994) Yan et al. (1987) Hung and Mayer (1985) Liu et al. (1991) Liu et aL (1991) BCttiger et al. (1987) Liu et aL (1991) Liu et al. (1991) Liu et al. (1983) Yan et al. (1987) Liu et aL (1987) Zhang and Liu (1994) Zhang (1995) Yan et al. (1987) Brenier et al. (1987) Liu et al. (1991) BCttiger et al. (1987) van Rossum et al. (1983) Liu et al. (1983) Liu et al. (1983) Jin and Liu (1995) Zhang et al. (1993) Liu and Zhang (1994) Cheng et aL (1984) Liu et al. (1991) Liu et al. (1991) BCttiger et al. (1989) Cheng et al. (1984) Liu (1985)
Several glasses have been formed in many systems. Or, at least one glass has been formed in some systems.
B. X. Liu
210 Table 8.2 Continued AHf
System
Ref.
Remarks
Any < 0
Au-Cu Cu-Rh Mo-Nb Fe-W
Rauschenbach (1986) Peiner and Kopitzki (1988) Liu (1982) Goltz (1982)
No glass has so far been formed.
Cu-Ir Ti-Zr
Peiner and Kopitzki (1988) Liu (1982)
Ag-Cr Ag--Cu Ag-Mo Ag-Nb Ag-Ni Ag-Ta Ag-V Au-Co Au-Ir Au-Mo Au-Os Au-Ru Au-W Co-Cu Cu-Fe Cu-Mo Cu-Nb Cu-Ta Fe-Nd Hf-Nb Hf-Ta Mo-Y Nb-Ti Nb-Zr Ta-Ti Ta-Y Ta-Zr Ti-Y Y-Zr
Liu et al. (1987) Liu et al. (1986) Jin et al. (1985) Jin and Liu (1996) Liu et al. (1986) Andersen et al. (1990) Andersen et al. (1990) Tsaus et al. (1981 ) Peiner and Kopitzki (1988) Pan and Liu (1995) Peiner and Kopitzki (1988) Meissner et al. (1987) Meissner et al. (1987) Liu et al. (1987) Huang and Liu (1987) Chen and Liu (1997) Andersen et al. (1990) Tsaur (1980) Yan et al. (1987) Jin et al. (1995) Liu et al. (1996) Zhang et al. (1995) Chen and Liu (1988) Jin and Liu (1994) Chen et al. ( 1991 ) Zhang and Liu (1995) Jin and Liu (1997) Liu et al. (1991 ) Liu et al. (1998)
Several glasses have been formed in many systems. Or, at least one glass has been formed in some systems.
Ag-Co Ag-Fe Ag-Ir Ag-Os Ag-Rh Ag-Ru Ag-W Au-Fe Au-Ni Au-Rh Cr-Cu Cu-Os Fe-Yb
Liu et al. (1998) Andersen et al. (1990) Peiner and Kopitzki (1988) Peiner and Kopitzki (1988) Peiner and Kopitzki (1988) Peiner and Kopitzki (1988) Hiller et al. (1989) Tsaur et al. (1981 ) Tsaur et al. (1981) Peiner and Kopitzki (1988) Shin et al. (1983) Meissner et al. (1987) Liu et al. (1991)
No glass has has so far been formed.
AHf = 0
AHf = 0
AHf
> 0
AHf>O
Several glasses have been formed. No glass has so far been formed.
Ion-Mixing
211
18.1%
+40
*20
10.0% 8 1%
1
0 Bulk L..
,I,,
Y
20 Mo
State 9
40
J
,
60
Concentration
,,
&
80
Mo
(at%)
Figure 8.6.
Calculated free energy diagram of the Y-Mo system. The numbers of 8.1%, 10.0% and 18.1% are for the fractions of interfacial atoms versus the total atoms in the multilayers, respectively.
8.5.2.3 Nominal and intrinsic glass-forming ability From both practical and theoretical points of view, the GFA and its prediction are of vital importance. The originally defined GFA was only to classify the systems into two categories of glass-forming and non-glass-forming ones in a qualitative way and the classification mostly depended on glass preparation technique employed. However, by employing powerful techniques, e.g., IM, many new metallic glasses have been obtained in those systems considered as non-glass-forming ones by LMQ. Apparently, the definition or the physical meaning of GFA needs to be extended. Firstly, GFA of a system should reflect qualitatively whether a metallic glass can be formed or not in the system. As mentioned above, the GFA determined by more powerful techniques frequently possessing higher cooling rates is expected to be closer to the real ability of the system. In other words, from a physical point of view, there exists an intrinsic GFA reflecting the capability of the system itself to form a metallic glass and it does not depend on the technique employed to produce the glass. Furthermore, it is necessary to have a quantitative measure of GFA to compare the capabilities of glass formation among the glass-forming systems. The experimentally determined GFR, in the author's view, can be taken as a quantitative measure of GFA. Apparently, the larger the GFR, the greater the GFA. It is true that if one employs a technique with a higher cooling rate, one may obtain the glassy phase in a broader composition range than that observed by a technique with a lower cooling rate. In this sense, the experimentally observed GFR can only be considered as a nominal GFA. Obviously, the greater the GFR observed by a specific technique, the closer the GFR is to the intrinsic GFA. Consequently, to approach the intrinsic GFA of a system, one should
B. X. Liu
212
Table 8.3 Some examples of metallic glasses formed in binary metal systems with a negative heat of formation by ion-mixing No.
System
AHf (kJ/mol)
MPAR (%)
Glass-forming range
1 2 3 4 5
Zr-Ru Ni-Nb Co-Nb Nb-Fe Co-Mo
-86 -44 - 38 -23 -7
90 86 86 85 76
25-75 at.% Ru 20-85 at.% Nb 20-77 at.% Nb 20-78 at.% Fe 35-77 at.% Mo
Table 8.4 Some examples of metallic glasses formed in binary metal systems with a positive heat of formation by ion-mixing. (Total thickness of the multilayered films is 40-50 nm.) No.
System
AHf (kJ/mol)
MPAR (%)
Number of interfaces
Glass-forming range
1 2 3 4 5
Zr-Nb Ag-Mo Ti-Ta Y-Nb Y-Mo
+6 +56 +3 +44 +35
70 100 100 100 100
7 11 15 17 18
at 12 or 60 at. % Nb at 30 or 80 at.% Mo 74-89 at.% Ta 28-44 at.% Nb 20-80 at.% Mo
look for a technique capable of generating cooling rates as high as possible. At present, IM can be considered as the technique capable of revealing the m a x i m u m GFA a m o n g the currently available glass producing techniques. The next issue is how to correlate the n o m i n a l G F R observed by IM and the intrinsic GFA of the system. Based on IM study, a parameter, namely MPAR, has been proposed by the author to predict the m a x i m u m possible G F R of a system. According to the definition, M P A R = 1 0 0 % - ( o r + 13), where c~ and/~ stand for the m a x i m u m solid solubilities at the two metal ends (Liu et al. 1987, Liu 1987). As IM can frequently form some supersaturated solid solutions, which would actually reduce the total width of the two-phase regions (Liu 1986), an empirical formula is therefore suggested for correlating the possible G F R and MPAR, i.e. G F R = 0.7MPAR, and it is reasonably compatible with the IM results in either positive or negative A H f systems, as shown by those examples listed in Tables 8.3 and 8.4. W h e n MPAR=0, corresponding to a system with a continuous solid solution phase diagram, G F R = 0 and no glass has so far been obtained in such systems, which are therefore considered as non-glass-forming ones.
213
Ion-Mixing
8.6. FORMATION OF METASTABLE CRYSTALLINE ALLOYS
8.6.1 Structural classification of the metastable crystalline phases Up-to-now, a number of metastable crystalline phases (MX) have been formed by ionmixing in both negative and positive A H i systems and they are classified into 5 structural categories (Liu et al. 1995), i.e., supersaturated solid solutions, the h.c.p.-I phases in h.c.p.- or f.c.c.-based alloys, the f.c.c.-I phases in h.c.p.-based alloys, the h.c.p.-II and f.c.c.-II phases in b.c.c.-based alloys. 8.6.1.1 Solid solutions
Ion-mixing of multilayered films can extend the solid solubility of the binary metal alloys resulting in the formation of a supersaturated solid solution always having the same simple crystal structure as the major alloying metal. Most of these alloy phases are identified as solid solutions. Their lattice constants agreed quite well with those predicted by Vegard's law, and they are therefore substitutional solid solutions (Liu and Jin 1987).
8.6.1.2 h.e.p.-I and f.e.e.-I phases A metastable crystalline phase with a similar h.c.p, structure has been formed by ionmixing in five binary (A-B) metal systems (Co-Au, Ti-Au, Co-Mo, Ni-Mo, and Ni-Nb) (see Table 8.5) where A refers to the first entry in the alloy designation. The MX phases were formed in A-rich multilayered films with an overall composition near A3B. These MX phases were named as h.c.p.-I, since they have all about the same spacing (dcpp) of the close packed planes (Table 8.5). Their formation was attributed to the valence electron effect. For a close-packed hexagonal structure, the maximum number of electron states per atom, n, in the Jones zone (including doubling for spin) can be calculated by the following equation (Jones 1960):
n=2-~
(a)2[ 1(a)2] c
1-~
c
'
(8.11)
where a and c are the lattice constants of the h.c.p, structure. According to this calculation, the n values of these five phases are almost identical, i.e., n -- 1.73 to 1.74, which is very close to the value of 7/4 = 1.75, corresponding to the well-defined Hume-Rothery 7/4 electron compound. These A3B alloy phases can therefore be considered as metastable electron compounds (Liu 1983). Table 8.6 shows that MX f.c.c.-I phases have formed in some h.c.p.-metal-ricia multilayered films by ion-mixing. As both h.c.p, and f.c.c, structures are built up by the stacking of close-packed atomic planes differing only slightly in the stacking order, one can easily transform into another. For example, sliding of atoms on the planes parallel to the (0002)hcp plane along (1 i00)hcp directions by a vector of 1/3(1 i00)hcp can transform the h.c.p, structure into f.c.c., and the reverse procedure transforms the f.c.c, structure to h.c.p (Liu and Zhang 1994a).
B. X. L i u
214
T a b l e 8.5 Structure data of the h.c.p.-I phase formed in five binary systems by ion-mixing. These phases were considered as metastable electron compounds. M X phase System Co-Au
Constituent metals
lattice constants
dcpp
lattice constants
dcpp
(nm)
(nm)
(nm)
(nm)
h.c.p,
d002 = 0.260
Co h.c.p,
d002 = 0.2023
a = 0.324, c = 0.521
a = 0.254, c = 0.407
c/a = 1.61
c/a = 1.59
Ref. Liu and Nicolet (1983)
Au f.c.c.
Ti-Au
h.c.p,
d002 = 0.263
a = 0.408
dl 11 = 0.2355
Ti h.c.p,
d002 = 0.234
a = 0.327, c = 0.526
a = 0.259, c = 0.469
c/a = 1.61
c/a = 1.59
Liu et al. (1982)
Au f.c.c.
Co-Mo
h.c.p,
d002 = 0.261
a = 0.408
d l l l = 0.2355
Co h.c.p,
d002 = 0.2023
a = 0.327, c = 0.522
a = 0.254, c = 0.407
c/a = 1.60
c/a = 1.59
Liu et al. (1983b)
M o b.c.c.
Ni-Mo
h.c.p,
d002 = 0.261
a = 0.3147
d110 = 0.2225
Ni f.c.c,
d l l l = 0.2034
a = 0.325, c = 0.521
a = 0.352
c/a = 1.60
Mo b.c.c. a = 0.3147
Ni-Nb
h.c.p, a = 0.328, c = 0.522
d002 = 0.261
Liu et al. (1983b)
d110 = 0.2225
Ni f.c.c. a = 0.352
d l l I = 0.2034
Kung et al. (1983)
Nb b.c.c.
c/a = 1.59
a = 0.3307
d110 = 0.233
Table 8.6 f.c.c.-I M X phases formed in Co-, Y-, Zr-, and Hf-rich multilayered films by room temperature 200 keV xenon ion mixing. System and composition
Mixing dose
Lattice parameter (nm)
Co94Mo6 ~ Co82Mo18
3 x 1015 X e + / c m 2
0.355 ~ 0.360
Co98Ta2 ~ Co80Ta20
3 • 1015 X e + / c m 2
0.355 ~ 0.365
Y73Mo27 ~ Y83MoI7
7 • 1015 X e + / c m 2
0.510
Y85TaI5 ~ Y75Ta25
1 x 1015 X e + / c m 2
0.518
Y25Zr75 ~ Y85ZrI5
3 • 1015 X e + / c m 2
0.467
Zr68Nb32 ~ Zr88Nb12
7 • 1014 X e + / c m 2
0.480
Hf75Nb25 ~ H f 9 0 N b l 0
5 • 1014 X e + / c m 2
0.455
Ion-Mixing
215
Table 8.7 h.c.p.-II and f.c.c.-II MX phases formed in Nb-, Ta-, Mo-rich multilayered films by room temperature 200 keV xenon ion-mixing. A-B
f.c.c.-b.c.c.
h.c.p.-b.c.c.
b.c.c.-b.c.c.
System and composition
h.c.p.-II
f.c.c.-II
Ni23Mo77 Ni20Nbs0 Ni35Ta65 Ag30Mo70 Ag33Nb67 Co20Mo80 Co23Nb77 Co30Ta70 Y22Mo78 Y25Ta75 Zr7Mo93 Zrl9Nbs1 Zr20Ta80 Hf17Nb83 Hf30TaT0 Fe24Mo76 Fe20Nbs0 Fe20Ta80
Yes
Yes
Yes No No No Yes Yes No Yes Yes Yes Yes No No No Yes
Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes
Yes
Yes
No
Yes
8.6.1.3 h.c.p.-H and f.c.c.-H phases based on b.c.c, metals In some Nb-, Ta-, and Mo-rich multilayered films, two MX phases, labeled h.c.p.-II and f.c.c.-II, were formed by ion-mixing, and their compositions were around AB3-4, in which B stands for the b.c.c metal. (see Table 8.7). Some of the systems have negative AHf, while combining with metals like Fe, Co and Ni and some have positive AHf when alloying with Ag, Y, Zr and Hr. Carefully performed experiments revealed that the h.c.p.II phase was observed as an intermediate phase in the process of forming the f.c.c.-II phase. Thus a two-step transition of b.c.c. --~ h.c.p. -+ f.c.c, was proposed and named as reverse martensitic transformation. First b.c.c, phase transforms to a h.c.p structure through a shearing mechanism. During shearing, the (111)bcc acted as the habit plane, while [ l l l]bcc (or [2110]hcp) acted as the shearing axis. The second step is to transform the h.c.p, phase to the f.c.c, phase. The shear can be achieved by sliding the atoms on planes parallel to the (0002)hcp along (1100)hcp directions by a vector of 1/3 (l l00)hcp. Although the relaxation period of ion-mixing is very limited, it is thought that it could meet the requirements of a shearing mechanism. The lattice parameters of h.c.p.-II and f.c.c.-II phases were deduced according to the two-step transition and these are in good agreement with the experimental results (Zhang 1995, Liu and Zhang 1992a).
B. X. Liu
216
8.6.2 Free energy calculation of the MX phases A metastable crystalline phase is of ordered structure, and so its entropy term can be neglected. The free energy of the MX phase is therefore A GMX = A H~I X -~- A Hi,ix -~- A H ~ x .
(8.12)
The expression of the chemical term is as follows A HI,cix = A HampX A I/" " A2/3 f A B ,
(8 13)
in which VA is the atomic volume of component A in the alloy, A Hamp is an amplitude concerning the magnitude of the electron redistribution interaction and is a constant for one specific metal system, faB is a function which accounts for the degree to which atoms of type A are surrounded by atoms of type B. In the calculation of faB, an empirical constant y is used to describe the short-range-order difference of solid solutions, amorphous phases, and ordered compounds, for which it is usually taken to be 0, 5, and 8, respectively. As an MX phase, except the solid solutions, is commonly considered as compoundlike, the value of y is therefore taken as 8 in the calculation. The elastic term of the MX phase can be derived from the equation A H ~ I X - - X A X B [ X A A H Me X ( B i n A )
where
-'F X B A
e , HMX(AinB)]
(8 . 14)
is the elastic contribution to the heat of solution of i in j, which can be calculated from the elastic constants of the constituent metals. The expression for the structural term is written as AOMx(iinj)e
A H ~ x = E(Z)
-
xAE(ZA)
--
xBE(ZB),
(8.15)
where E(Z), E(ZA) and E(ZB) are the structural stabilities of the alloy phase and the pure A and B metals, and Z, ZA and ZB are the mean number of valence electrons of the alloy phase and the number of valence electrons of pure metals A and B, respectively. For an MX phase at a composition in the vicinity of an intermetallic compound, the elastic term should be included. This term was from the difference in structures between the MX phase and that of the intermetallic compound. This is for the case that the MX phases were formed at A3B stoichiometry. Another situation is that the MX was formed at compositions near the eutectic point, where no equilibrium compound exists. In this case, the MX phase can be considered as ordered, so as the elastic term can be relaxed down to zero because of ordering. This situation corresponds to the formation of the MX phases with a composition of AB3. In short, with the above assumption, the formula for calculation for the A3B MX phase AGMx -- A G ~ x 4- A G ~ x + A G ~ x ,
(8.16)
and for the AB3 MX phase AGMx-
AG~x + AG~x.
(8.17)
Ion-Mixing
217
~" 12001-1 bcc(Mo) o.. I]
i) 91(P~
~ .
80(,[]
~
MoN
~- 6o0[/
i3
Y
4(t(t U
. . . . . . . . . 2(t 40 6(t
Mo
8()
Ni
(a)
~
"~ +8 E
bcc solution
fcc solution
..x +4
=
0
r
hcp-ll ~
~,, -4
~kJ "~
fcc-ll --
Mo
(b)
,A ....
9
20
9
J
40
9
,~
60
,
9
9
84)
9
Ni
Ni concentration (at.%)
Figure 8.7. Equilibrium phase diagram of the Ni-Mo system; (b) Calculated free energy diagram of the Ni-Mo system.
The free energy curves of the MX phases formed in the Ni-Mo system are calculated, using the above equations, as an example. Figure 8.7(a) presents the equilibrium phase diagram. Figure 8.7(b) shows the constructed free-energy diagram of the system, in which the free-energy curves of all the concerned phases are included. One sees that the diagram is a little complicated because some newly formed MX phases are also included. Before employing the diagram to interpret the alloy phase formation, an important question to be considered is if the calculated results are relevant, since the model used in the calculations is still semi-quantitative. The author has proposed an alternative way for checking the relevance of the calculation, i.e., the steady-state thermal annealing of the multilayered films at some important compositions. The idea was to see whether the phase evolution in
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B. X. Liu
the films upon thermal annealing matches with that deduced from the calculation. Let us consider two examples in which the phase appearance in the samples is as follows:
300 ~ 450 ~ 650 ~ 800 ~ Niv5Mo25 30min > Am~ min > hcp-I20 min > Ni + hcp-II30 mi n > Ni + fcc-II
350~ 500~ 600~ Ni25Mo75~ > Amor.~ > hcp-II~ > fcc-II 60 min 30 min 30 min Apparently, the phase evolution sequences of these two reactions are in good agreement with that deduced from the calculated free energy diagram, thus confirming the applicability of the Miedema's model in this respect (Zhang and Liu 1994a). The author's group has also performed similar annealing experiments for several positive A H f systems and these results are also in support of the calculations (Chen and Liu 1996). The use of calculated free energy diagrams can explain the alloy phase formation upon IM, as discussed in the previous sections by some representative examples.
8.7. I N T E R F A C E - G E N E R A T E D
S O L I D - S T A T E V I T R I F I C A T I O N IN S Y S T E M S W I T H A
P O S I T I V E H E A T OF F O R M A T I O N
In 1983, Schwarz and Johnson reported for the first time that solid-state reaction (SSR) of multilayers consisting of two different crystalline metals (La and Au) can lead to the formation of an amorphous alloy. They argued that spontaneous vitrification was expected to take place in metal systems with a negative A H i and one metal diffusing anomalously fast in the other. However, the author's group has found very recently that in some binary metal systems with positive A H f , vitrification could be achieved by SSR, e.g., in the HfTa (Liu et al. 1996), Y-refractory metals (Chen and Liu 1997) and Cu-Ta (Pan et al. 1998) systems. Consider the Y-Nb system as an example, Figure 8.8 shows the X-ray diffraction patterns of the Yv2Nb28 multilayered films before and after thermal annealing. It should be emphasized that similar solid-state vitrification was also observed by author's group very recently in some other systems with a positive A H f (Chen and Liu 1997a). To test whether a large size difference is a decisive prerequisite in solid-state amorphization, the Au-Ta system with a negative A H f o f - 4 9 kJ/mol (Pan et al. 1995) was selected as the atomic radius ratio and the volume ratio between Au and Ta are 1.008 and 0.936, respectively. When the Auz3Ta77 multilayers were heated up continuously at a rate of 5 K/min to 200~ partial amorphization was observed. After maintaining at 200~ for 30 min, the annealing temperature was raised to 350~ and kept for 30 min. The structure of the multilayers turned entirely into an amorphous state. When the temperature was further raised to about 650~ and kept for 20 min, diffraction lines from metals Au and Ta emerged again, indicating crystallization of the Auz3Ta77 amorphous alloy. These results suggest that solid-state vitrification can also be achieved in a system with a small atomic size difference.
Ion-Mixing
219
J
.1500
(a) .
.
.
.
0i 0
.,,.i
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.e
(b)
~0
i
30
|
I
40 50 20 (degree)
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-
Figure 8.8. X-ray diffraction patterns of the Y72Nb28 multilayered films with 121 layers and a total thickness of 350 nm before and after thermal annealing: (a) as-deposited; (b) amorphized by annealing at 300~ for 1.5 hours.
8.8. C O N C L U D I N G R E M A R K S
(1) In the past 15 years, study of IM of multilayered films scheme has provided extensive data, which promoted significantly the non-equilibrium processing of materials. IM has so far synthesized a great number of new amorphous as well as new MX alloys. It is anticipated that this technique is capable of synthesizing more new metallic glasses in those not yet studied systems known as non-glass-forming ones. (2) Miedema's theory was employed for free energy calculation and steady-state thermal annealing results showed that the calculated energetic sequence of the related phases was relevant, at least in its outline, suggesting the applicability of Miedema's theory in this respect. (3) Taking into account the interfacial free energy, a unified thermodynamic model was proposed for predicting the possibility of metallic glass formation by IM in binary metal systems with either negative or positive AHz, and for a reasonable explanation for the formation of MX phases. (4) The interface concept was also tested by solid-state reaction experiments in some positive A Hf systems and it turned out that metallic glasses were indeed obtained by thermal annealing in the properly designed multilayered films with enough fraction of interfacial atoms. These results offer a neat evidence demonstrating the decisive role of the interfacial free energy in amorphization taking place in the multilayered films.
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B. X. Liu
(5) Kinetically, due to the very restricted conditions available in IM for a crystalline phase to nucleate and grow, only simple-structured phases can be formed and they feature a close crystallographic correlation with their matrix. (6) Based on the IM and SSR studies, a new glass-producing technique is developed, i.e., the multilayer technique, in which the multilayered films are firstly designed according to the interfacial free energy calculation and then either by thermal annealing or ion irradiation. (7) Further studies are still needed in this area, especially in the systems with a positive A Hf. For instance, detailed growth of the amorphous phase, local microstructures and the abnormal properties of the new metallic glasses formed in such systems are highly interesting topics to be investigated by IM as well as SSR technique.
ACKNOWLEDGMENTS
The author is grateful to his current and previous students for their continuous contributions to this research project, especially to Mr. Y.G. Chen for his substantial assistance in writing and finalizing this chapter. Special thanks go to the Editor of this book, Professor C. Suryanarayana, for his critical reading of the manuscript of this chapter and for many helpful suggestions. The continuous financial support from National Natural Science Foundation of China under several projects and by the Administration of Tsinghua University are greatly appreciated.
REFERENCES Alonso, J. A. and Simozar, S. (1983) Solid State Comm., 48, 765. Alonso, J. A., Gallego, L. J. and Simozar, S. (1990) Nuovo Cimento, D12, 587. Andersen, U. A., BCttiger, J. and Dyrbye, K. (1990) Nucl. Instru. Meth. Phys. Res., B51, 125. Bcttiger, J., Pampus, K. and Torp, B. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 696. Bcttiger, J., Dyrbye, K., Pampus, K. and Poulsen, R. (1989) Phil Mag., A59, 569. Brenier, R., Thevenard, P., Capra, T., Perez, A., Treilleux, M., Romana, L., Dupuy, J. and Brunel, M. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 691. Carter, G. and Grant, W. A. (1976) Ion Implantation of Semiconductors. (Edward Arnold Ltd., London). Chen, Y. G. and Liu, B. X. (1996) Appl. Phys. Lett., 68, 3096. Chen, Y. G. and Liu, B. X. (1997a) J. Alloys and Compounds, 261, 217. Chen, Y. G. and Liu, B. X. (1997b) Nucl. Instr. Meth. in Phys. Res., B127/128, 145. Chen, Y. G. and Liu, B. X. (1997c) J. Appl. Phys., 82, 3815. Chen, Y. G. and Liu, B. X. (1988) Mater Sci. Eng., B ( to be published). Chen, Y. G., Zhang, Q. and Liu, B. X. (1997) Nucl. Instr. Methods Phys. Res., B124, 523. Cheng, G. A. (1986) Thesis, Tsinghua University, Beijing, China. Cheng, Y. T., Johnson, W. L. and Nicolet, M.-A. (1984) in Proceedings of MRS Fall Meeting, Boston, MA. Chu, W. K., Mayer, J. W. and Nicolet, M.-A. (1978) Backscattering Spectrometry. (Academic Press, New York). Clemens, B. M. (1986) Phys. Rev., B33, 7615. Davies, H. A. (1983) in Amorphous Metallic Alloys, ed. Luborsky, EE. (Butterworths, London), p.13. de Boer, F. R., Boom, R., Miedema, A. R., Niessen, A. K., and Mattens, W. C. M. (1989) Cohesion in Metals: Transition Metal Alloys. (North-Holland, Amsterdam). Dearnaley, G., Freeman, J. H., Nelson, R. S., and Stephen, J. (1973) Ion Implantation. (North-Holland Publishing Company, Amsterdam). Ding, J. R., Che, D. Z., Zhang, H. B., Tao, K. and Liu, B. X. (1992) Appl. Phys. Lett., 60, 944. Gerkema, J. and Miedema, A. R. (1983) Surf. Sci., 124, 351. Giessen, B. C. and Whang, S. H. (1980) J. Physique, C8, 95.
Ion-Mixing
221
Goltz, G., Fernandez, R., Nicolet, M.-A. and Sadana, D. K. (1982) in Metstable Materials Formation by Ion Implantation. eds. Picraux, S. T. and Choyke, W. J. (North-Holland, Amsterdam), p. 227. Hafner, J. (1981) in Metallic Glasses, eds. Gtintherodt, H. J. and Beck, H. (Springer-Verlag, Berlin), p. 91. Hiller, W., Buchgeister, M., Eitner, P., Kopitzki, K., Lilienthal, V. and Peiner, E. (1989) Mater Sci. Eng., All5, 151. Huang, L. J. and Liu, B. X. (1987) Nucl. Instr. Methods Phys. Res., B18, 256. Hung, L. S. and Mayer, J. W. (1985) Nucl. Instr. Methods Phys. Res., B7/8, 657. Hung, L. S., Nastasi, M., Gyulai, J. and Mayer, J. W. (1983) Appl. Phys. Lett., 42, 672. Jin, O. and Liu, B. X. (1994) J. Phys.: Condensed Matter, 6, L39. Jin, O. and Liu, B. X. (1995) J. Phys.: Condens. Matter, 7, 2667. Jin, O. and Liu, B. X. (1996) Nucl. Instr. Meth. Phys. Res., Bl14, 56. Jin, O. and Liu, B. X. (1997) J. Non-Cryst. Solids, 211, 180. Jin, O., Ye, Y. and Liu, B. X. (1995a) J. Phys. C: Condens. Matter, 7, 7959. Jin, O., Zhang, Z. J. and Liu, B. X. (1995b) Appl. Phys. Lett., 67, 1524. Jones, H. (1960) in The Theory of Brillouin Zones and Electronic States in Crystals. (North-Holland Publ. Co., Amsterdam), p. 191. Klement, W., Willens, R. H. and Duwez, P. (1960) Nature, 187, 869. Knapp, J. A. and Follstaedt, D. M. (1987) Nucl. Instr. Methods Phys. Res. B19/20, 611. Kung, K. T.-Y., Liu, B. X. and Nicolet, M.-A. (1983) Phys. Stat. Sol., (a) 77, 355. Lin, C. J. and Spaepen, E (1984) in Rapid Solidification of Metastable Materials, eds. Kear, B. H. and Giessen, B. C. (North-Holland Publishing Co., Amsterdam). Liu, B. X. (1982) California Institute of Technology, Pasadena, CA, unpublished research. Liu, B. X. (1983) Phys. Stat. Sol., (a) 75, K77. Liu, B. X. (1985) Nucl. Instr. Methods Phys. Res., B7/8, 547. Liu, B. X. (1986) Phys. Stat. Sol., (a) 94, 11. Liu, B. X. (1987) Mater. Lett., 5, 322. Liu, B. X. and Jin, O. (1997) Phys. Stat. Sol., (a)161, 3. Liu, B. X. and Nicolet, M.-A. (1983) Thin Solid Films, 101, 201. Liu, B. X. and Zhang, Z. J. (1994a) J. Mater. Res., 9, 357. Liu, B. X. and Zhang, Z. J. (1994b) Phys. Rev., B49, 12519. Liu, B. X., Nicolet, M.-A. and Lau, S. S. (1982) Phys. Stat. Sol., (a) 73, 183. Liu, B. X., Johnson, W. L., Nicolet, M.-A. and Lau, S. S. (1983a) Appl. Phys. Lett., 42, 45. Liu, B. X., Johnson, W. L., Nicolet, M.-A. and Lau, S. S. (1983b) Nucl. Instr. Meth., 209/210, 229. Liu, B. X., Clemens, B. M., Gaboriaud, R., Johnson, W. L. and Nicolet, M.-A. (1983) Appl. Phys. Letters, 42, 624. Liu, B. X., Huang, L. J., Li, J. and Ma, E. (1986) in Thin Films-Interfaces and Phenomena, eds. Nemanich, R. J., Ho, P. S. and Lau, S. S. (North-Holland, Amsterdam), p. 215. Liu, B. X., Ma, E., Li, J. and Huang, L. J. (1987a) Nucl. Instr. Methods. Phys. Res., B19/20, 682. Liu, B. X., Ma, E., Li, J., and Huang, L. J. (1987b) NucL Instr. Methods. Phys. Res., B19/20, 682. Liu, B. X., Ding, J. R., Zhang, H. B. and Che, D. Z. (1991) Phys. Stat. Sol., (a) 128, 103. Liu, B. X., Che, D. Z., Zhang, Z .J., Lai, S. L. and Ding, J. R. (1991) Phys. Stat. Sol., (a) 128, 345. Liu, B. X., Zhang, Z. J., Jin. O. and Pan, E (1995) Nucl. Instr. Methods Phys. Res., B106, 17. Liu, B. X., Jin, O. and Ye, Y. (1996) J. Phys. C: Conden. Matter, 8, L79. Liu, B. X., Zhang, Z. J., and Jin, O. (1998) Mater. Lett. (submitted). Lopez, J. M. and Alonso, J. A. (1985) Z. Naturforsch., A40, 1199. Mayer, J. W. and Lau, S. S. (1988) Electronic Materials Science: For Intergrated Circuits in Si and GaAs. (Macmillan Publishing Company, New York). Mayer, J. W., Tsaur, B. Y., Lau, S. S. and Hung, L. S. (1981) Nucl. Instr. Methods, 182/183, 1. Meissner, J., Kopitzki, K., Mertler, G. and Peiner, E. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 669. Miedema, A. R. (1976) Philips Tech. Rev., 36, 217. Miedema, A. R. (1978) Z. Metallkde., 69, 287 and 455. Murr, L. E. (1975) Interfacial Phenomena in Metals and Alloys. (Addison-Wesley, London). Niessen, A. K. and Miedema, A. R. (1983) Ber. Bunsenges. Phys. Chem., 87, 717. Niessen, A. K., Miedema, A. R., de Boer, E R. and Boom, R. (1988) Physica, B151, 401. Ossi, P. M. (1990) Phys. Star. Sol., (a) 119, 463. Pan, E and Liu, B. X. (1995) J. Phys.: Condens. Matter, 8, 383.
222
B. X. L i u
Pan, E et al., (1998) unpublished data. Pan, E, Chen, Y. G. and Liu, B. X. (1995) AppL Phys. Lett., 67, 780. Pan, E, Chen, Y. G., Zhang, Z. J., and Liu, B. X. (1996) J. Non-Cryst. Sol., 194, 305. Peiner, E. and Kopitzki, K. (1988) Nucl. Instrum. Methods Phys. Res., B34, 173. Rauschenbach, B. (1986) Appl. Phys., A40, 47. Rauschenbach, B. and Hohmuth, K. (1987) Nucl. Instr. Methods Phys. Res., B23, 323. Schmid, A. and Ziemann, P. (1985) Nucl. Instr. Methods Phys. Res., B7/8, 581. Schwarz, R. B. and Johnson, W. L. (1983) Phys. Rev. Lett., 51, 415. Shin, S. M., Ray, M. S., Rigsbee, J. M. and Greene, J. E. (1983) Appl. Phys. Lett., 43, 249. Takayama, S. (1976) J. Mater. Sci., 11, 164. Thompson, M. W. (1969) Defects and Radiation Damage in Metals. (Cambridge University Press, Cambridge). Townsend, E D., Kelly, J. C. and Hartley, N. E. W. (1976) Ion Implantation, Sputtering and their Applications. (Academic Press, London), p. 31. Tsaur, B. Y. (1980) Thesis, California Institute of Technology, Pasadena, CA. Tsaur, B. Y., Lau, S. S., Hung, L. S. and Mayer, J. W. (1981) Nucl. Instr. Methods, 182/183, 67. Turnbull, D. (1955) Impurities and Imperfections. (Amer. Soc. Met., Cleveland), p. 121. van Rossum, M., Shreter, U., Johnson, W. L. and Nicolet, M.-A. (1983) in Proceedings of MRS Fall Meeting, Boston, MA. von Allmen, M. (1983) in Laser-Beam Interaction with Materials, ed. Carl, J. (Washington International Society of Optical Engineering, Washington), p. 140. Yan, Z. H., Liu, B. X., and Li, H. D. (1986) Phys. Stat. Sol., (a) 94, 483. Yan, Z. H., Liu, B. X., and Li, H. D. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 700. Zhang, Z. J. (1995) Ph.D. Thesis, Tsinghua University, Beijing, China. Zhang, Z. J. and Liu, B. X. (1994a) J. Appl. Phys. 76, 3351. Zhang, Z. J. and Liu, B. X. (1994b) J. Appl. Phys., 75, 4948. Zhang, Z. J. and Liu, B. X. (1994c) J. Phys.: Condens. Matter, 6, 2647. Zhang, Z. J. and Liu, B. X. (1994d) J. Phys.: Condens. Matter, 6, 9065. Zhang, Z.J. and Liu, B. X. (1995) Phys. Rev., B51, 16475. Zhang, Z. J., Bai, H. Y., Qiu, Q. L., Yang, T., Tao, K. and Liu, B. X. (1993) J. Appl. Phys., 73, 1702. Zhang, Z. J., Jin, O. and Liu, B. X. (1995) Phys. Rev., B 51, 8076.
Chapter 9 Physical Vapor Deposition 9.1. Introduction 9.2. Development of PVD 9.3. Deposition Methods 9.3.1 Evaporation 9.3.2 Molecular Beam Epitaxy (MBE) 9.3.3 Sputtering 9.3.4 Ion-Assisted Deposition 9.3.5 Magnetron Sputtering 9.4. Influence of Energy on Coatings 9.4.1 Film Morphology and Density 9.4.2 Nucleation and Adhesion 9.4.3 Metastable Phases 9.4.4 Microstructure 9.5. Applications of PVD Coatings 9.5.1 Optical Coatings 9.5.2 Corrosion Protective Coatings 9.5.3 Hard Coatings 9.6. Future Trends Acknowledgment References
225 225 228 228 229 229 231 235 238 238 239 240 240 243 243 245 246 248 249 249
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Chapter 9 Physical Vapor Deposition JOHN S. COLLIGON
9.1. INTRODUCTION Physical Vapor Deposition (PVD) is the general name given to coating processes where the transport of material to the substrate is effected by a physical driving mechanism. Such mechanisms include evaporation, sputtering, ion-plating, and ion-assisted sputtering. Often chemical reactions do take place to form coatings of oxides, nitrides and carbides, but the transport process does not generally rely on chemical reactions. The coated samples can usually be maintained at a temperature between ambient and 100~ during the deposition and the working pressure is below 10 -2 Pa. PVD thus differs quite markedly from Chemical Vapor Deposition (CVD) which occurs at atmospheric pressure, or just below, and at temperatures normally in excess of 500~ CVD also requires the use of complex gaseous species (such as fluorides, organometallics) which need to be handled very carefully whereas PVD requires either no gas, for evaporation, or, relatively simple species (Oxygen, Nitrogen) for most of the other PVD processes. Recent improvements in PVD techniques, in particular the use of plasmas and the unbalanced magnetron for the deposition process, have led to its wide application in industry. Such coatings are used for corrosion protection, to provide hard surfaces for engineering applications, to form decorative coatings on a range of household and industrial products, for energy-saving window glass, for media storage systems such as compact discs and for a range of microelectronics applications. Table 9.1 lists some of the conference series that cover this huge field. The present chapter starts with a brief historical review of the subject. It then describes the main PVD methods and continues with a discussion of the relation between deposition conditions and film properties. Some of the main applications of PVD coatings are then described and the chapter concludes with an outline of some of the possible future trends in this important industrial field.
9.2. D E V E L O P M E N T
OF P V D
The first recorded mention of a PVD process was by Grove (1852) in 1852 who observed the formation of a deposit in a glass tube in which a discharge had been running and identified this as being the same material as one of the electrodes; this was the first reported evidence of the ejection of atoms from an ion-bombarded material, now known as sput225
226
J. S. Colligon
Table 9.1. Established conference series containing papers on PVD. 9 9 9 9 9 9 9 9 9 9 9 9 9 9
InternationalConference on Metallurgical Coatings and Thin Films (ICMCTF) AnnualTechnical Conference of the Society of Vacuum Coaters E-MRSSpring Meeting InternationalConference on Crystal Growth and International Conference on Vapor Growth and Epitaxy InternationalVacuum Congress and International Conference on Solid Surfaces InternationalConference on Thin Films InternationalConference on Nanometer-scale Science and Technology (Nano) InternationalConference on Ion Beam Modification of Materials (IBMM) InternationalConference on Surface Modification of Metals by Ion Beams (SMMIB) InternationalConference on Advances in Surface Engineering EuropeanConference on Diamond, Diamond-like Materials, Nitrides and Silicon Carbide (Diamond) InternationalConference on Plasma Surface Engineering (PSE) NationalSymposiumof the American Vacuum Society MRS Fall Meeting
tering and one of the main PVD processes. A few years later Faraday (1857) exploded metal wires in vacuum and thereby produced an evaporated film. Many years then passed before the concept of ionising some of the depositing species was proposed by Berghaus (1938) and, subsequently exploited by Mattox (1963) in the process known as ion-plating. Ion-plated coatings were seen to have superior adhesion properties which were identified with the additional energy imparted to some of the evaporated atoms which had been ionized and to a general increase in energy of support gas and other species bombarding the growing film. The process operated at a relatively high pressure (0.5 Pa) however, so that many collisions occurred in the gas phase which tended to reduce the energies of ionized species and increase those of neutral particles. The many collisions allowed depositing species to be bounced around shaped substrates to the rear side, giving the technique so-called high throwing power, but the parameters were not easy to control. Weissmantel (1976) proposed the use of ion beams in the deposition process where material was either evaporated onto a surface being simultaneously bombarded by ions, or was sputtered by one ion beam and the deposited film bombarded by a second beam. Here the precise control of the added energy and ion to atom arrival ratio could be achieved. A review of the early work was given by Weissmantel et al. (1979). This type of system, further exploited by Harper et al. (1982a, 1982b, 1983), Sainty et al. (1984), Martin et al. (1983, 1984), Colligon (1995), Kheyrandish et al. (1994) and many others (Rossnagel and Cuomo 1988), has allowed controlled studies of the influence of added energy on film composition, stress, density, microstructure, hardness, wear and durability to be made. An alternative approach has been proposed by Yamada and Takagi (1995) who developed an ion cluster-beam method. In this method, illustrated schematically in Fig. 9.1, a beam of clusters of atoms from 2 - 2 0 in each cluster is created by expanding a vapor of the metal through a nozzle. Some atoms in the clusters are then ionized by an electron beam so that the whole cluster may be accelerated toward the negative substrate. How-
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ce material Figure 9.1. Schematic diagram of ionized cluster-beam system
ever, all the above ion beam systems have a common limitation in that they do not allow the production of coatings over large areas (normally a maximum of a few square centimeters) nor do they allow high growth rates to be attained (normally maximum rates are 0.2-1 nm/s). There are some developments in the cluster-beam method by Gspann (1993) who has recently produced cluster-beam sources with much improved deposition rates but the treated areas are still of order of a few square centimeters. The major breakthrough in industrial processing has however come from the new generation of unbalanced magnetron systems (Window and Savvides 1986a, b, Savvides and Window 1986) which have, in the 1990s, replaced many simple evaporation methods and ion-plating techniques which had been the main coating methods during the period from the 1960's. Film deposition rates of 10-500 nm/s are now possible and the sample area can be several square meters. Similar rates of deposition, but not such large area capability, have been found for recent plasma-based systems, including the pulsed plasma methods (Conrad et al. 1990). All these systems can be developed to produce reliable coatings and, in some cases, can produce metastable phases of alloy materials because the energetic atoms are suddenly trapped into rapidly cooling surroundings. It is not always possible to relate the film properties to the deposition parameters because there are complex interactions between magnetron power, system pressure, ion/atom arrival ratio and deposition energies. Fundamental studies on new coatings are, therefore, often best undertaken on ion-beam systems.
J. S. Colligon
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9.3. D E P O S I T I O N M E T H O D S
This review will outline the main features of the various PVD methods but will concentrate more on the more recent ion-assisted deposition (IAD) techniques where additional energy is provided to the growing film by ion-bombardment. Earlier studies of traditional evaporation and sputtering methods have been fully reviewed in several textbooks by Holland (1956), Chopra (1969), Vossen and Kern (1978) and Bunshah (1982).
9.3.1 Evaporation In this process materials are vaporised by heating using direct resistance-heating, radiation, eddy current heating, electron-beam heating, or running an arc discharge. One advantage of the technique is that the evaporation can be carried out in a UHV chamber so that contamination of the substrate by residual gases is minimal. For best purity an electron-beam heating system should be used as the material can then be evaporated from the center of the crucible in which it is held and reactions with the crucible walls are minimised. A schematic diagram of a typical e-beam coating system is shown in Fig. 9.2. Fundamental equations for evaporation rates and tables of evaporation temperatures and procedures are given in Chapter 4 of the book edited by Bunshah (1982). Some decomposition of chemical compounds always occurs and special procedures have to be followed if fully stoichiometric films are to be formed. The energies of species arriving at the substrate are only of the order of 0.05-0.1 eV which means they have limited energy for surface migration to find good nucleation sites. Substrates must be free of contaminants to allow reasonable levels of film adhesion to be attained so considerable attention has to be paid to cleaning procedure. Because of the lack of surface mobility
Physical Vapor Deposition
229
there is a high probability that the coating will contain voids and will be less dense than the bulk material. This leads to the formation of stress in the coating, usually tensile, which often results in the film delaminating and curling up.
9.3.2 Molecular beam epitaxy (MBE) A special technique of evaporation has been developed, primarily for the semiconductor industry, in which component species are injected into a UHV chamber from gas effusion cells and form the desired film on a substrate. This is known as Molecular Beam Epitaxy. In some systems organometallic gases are used which react at the high temperature substrate to form the desired deposit and a residual volatile species. Following the work of Giinther (1958) an alternative three-temperature method was developed which allowed III-V semiconductor compounds to be deposited from the pure constituent elements; usually the type-III element cell was run at a higher temperature than that for the more volatile type-V element and the substrate temperature was chosen to be high enough to re-evaporate excess type-V species. A review of fundamental aspects of MBE has been published by Foxon and Joyce (1981). In later work an important correlation between the oscillations in intensity of RHEED signals and layer coverage was discovered, allowing the number of layers deposited to be calibrated in situ (Neave et al. 1983). The whole field is well reviewed by Herman and Sitter (1989) and Chow (1991) and a comprehensive bibliography of early data relating to MBE of III-V compounds (from 1958-1983) has been published by Ploog and Graf (1984). Later improvements have allowed the deposition conditions to be improved by applying reactive gas in pulses (Goodman and Pessa 1986) which allows growth at lower temperatures and hence reduces diffusion to produce sharper boundaries. The choice of gas species and effusion cell temperature is very important as the structure of the molecule forming the film (e.g., dimer or trimer) is of major importance. The MBE process has been linked to other stages in the semiconductor production line as described, for example, by Migauchi and Hashimoto (1986) and by Melngailis (1987).
9.3.3 Sputtering When an ion strikes a material its energy is transferred by a series of collisions to the atoms of the first few layers of that material. Some of these atoms will move toward the surface and, if their energy is greater than the binding energy they will escape. The ratio of number of ejected atoms per incident ion is known as the sputtering coefficient and, for an incident ion of energy about 1 keV, it has a value for most materials ranging between 0.1 and 2 atoms per ion. A full review of the mechanisms and major trends in the sputtering phenomena was published by the author in 1968 (Carter and Colligon 1968, Chapter 7). The sputtering coefficient increases, after an initial quadratic behavior below 100 eV, almost linearly with ion energy as the ion range increases and more and more atoms participate in the energy exchange process. However, when the ion energy increases such that the ion range exceeds several atomic layers, the energy is too deep to be transferred back to the surface and the sputtering coefficient begins to decrease again.
230
J. S. Colligon
The fundamental understanding of the process has not changed significantly since the late 1960s but the theory of the process has been fully developed, the model by Sigmund (1969, 1971) proving to be very successful in predicting sputtering coefficients. Of particular importance to thin film deposition is the energy of the sputtered species which has a peaked distribution with a maximum at about 4-7 eV and a long tail down to over 50 eV. These energies are more than an order of magnitude greater than those of evaporated species and allow the depositing species more mobility on the surface to form denser, well-bonded films. One problem with the method is the change in composition of the surface of a sputtered compound material which arises from preferential sputtering of components. This results primarily as a result of the different energy transfer to different mass species. Fortunately, after a period of bombardment, the system is selfcompensating since the number of atoms of the species sputtering most readily is reduced and the rate of further sputtering of this component returns to the stoichiometric level. It is always important therefore to sputter a compound target for a period of time and to cover the substrate during this stabilizing period. Another problem from compound targets is that the spatial distributions of component ejected species can be different. Thus a substrate at a specific position in space facing the target may not collect component atoms in the correct stoichiometric proportions, even though the global total sputtering of species is stoichiometric. Clearly it is preferable, where possible, to move substrates so that they sweep through these spatial distribution variations. Apart from this problem, the process of sputtering is very versatile and control of the deposition rate is relatively simple. The first commercial systems were based on so-called planar diode sputtering and had the form illustrated in Fig. 9.3. Electrons are accelerated towards the anode and have ionizing collisions with gas atoms. These ions then bombard the cathode, causing sputtering and the liberation of further electrons which cause the discharge to be sustained. There are also many ion-gas collisions leading to charge exchange so that the target is bombarded by a mixture of ions and neutrals with energies below the total applied voltage. Typical operating conditions for dc metal sputtering are; cathode current density: 1 mA/cm 2, discharge voltage: 3 kV, argon pressure: 10 Pa, cathode-to-substrate distance: 40 mm, deposition rate: 40 nrn/min. For sputtering of insulating materials a radio-frequency (rf) supply has to be used, usually at the allowed 13.56 MHz frequency. These rf systems run at lower pressures, in the range 1-2 Pa. Although materials can be deposited using planar sputtering systems, the deposition rate is rather low, the ionization efficiency in the plasma is also low and the substrate becomes heated. These limitations have been overcome by the introduction of magnetron sputtering systems which are described in the next section. Deposition rates in these systems are as high as several microns per minute and can be controlled by varying the power density supplied to the target region, the target area, coating pressure and target-substrate distance. Although the substrates are near to the ionized plasma region the degree of ion bombardment of the growing film is low in these systems. The film quality is thus similar to that obtained for other sputtering methods but with a slight improvement in film density arising from the lower operating pressure; atoms sputtered from the target do not suffer many collisions in transit and therefore maintain, on average, a higher energy.
231
Physical Vapor Deposition
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9.3.4 Ion-assisted deposition
Evidence that additional ion bombardment of a surface being coated improved the adhesion and film properties was demonstrated in the early 1960s by Mattox (1963, 1965) using a system known as ion-plating, which was patented much earlier by Berghaus (1938). A typical ion-plating system is illustrated schematically in Fig. 9.4. The coating material is evaporated using an electron-beam and a discharge is struck in the upper region using a support gas such as argon or, for nitrides, nitrogen. Operating pressures are generally in the region 0.1-1 Pa which means that ions and atoms in the discharge make many collisions in passing to the negatively biased substrate. The substrate is thus bombarded by a variety of atomic, ionized and molecular species, including ions and atoms of the coating material with energies from thermal values up to the applied voltage of several keV. A further effect is the irradiation of the substrate and growing film by photons from the discharge. These energy-assisted deposition conditions have been shown to produce coatings with improved adhesion (Mattox 1965, Rigsbee et al. 1986, McLeod and Mah 1974, Enomoto and Matsubara 1975, Takagi et al. 1974, Teer 1983, Ahmed and Teer 1984), increased density (Mattox and Kominiak 1972), improved abrasion resistance (Lardon et al. 1978, Broszeit et al. 1985), and enhanced optical reflectivity (Henderson 1977, Varasi et al. 1985). However the film-growth parameters cannot be controlled independently; for example, a change in pressure also effects gas collisions and therefore energies of arriving particles.
232
J. S. Colligon
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An altemative plasma-based ion-assisted deposition system is the arc-plating system. Here a discharge is run in the abnormal glow region of the current-voltage characteristic for a diode system. Extremely high (many amperes) currents flow in the discharge which forms an arc with the current entering the target at a small (few mm dia) spot. Target material is vaporized and, in the intense discharge, a high fraction of this material is then ionized. A biased substrate will therefore collect both atoms and ions of the film species (together with neutral species). The arc is struck with an igniter (used to produce fieldemission of electrons) and starts with additional gas in the chamber (such as argon) but, the arc once struck, the argon can be removed and the discharge maintained in the vapor. In operation the arc moves randomly over the target surface but this can be controlled by applying a movable magnetic field (steered arc system). Film quality is as good as for ion-plating with the exception that there is a tendency for large droplets of target material to erupt from the surface and fall on the substrate, thus spoiling film uniformity. This problem can be minimized by running at lower arc currents (but with associated lower deposition rates) or by moving a magnet below the cathode to "steer" the arc (Ramalingam et al. 1987, Erttirk and Heuvel 1989). Later versions have however introduced an arc deflecting column to entrain the arc around a 90 degree bend so that the substrate may be
233
Physical Vapor Deposition
coit'~
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mounted out of the line of sight from the target. This is known as the filtered arc system and is shown in Fig. 9.5 (Vyskocil and Musil 1990). In this situation only the ionized fraction of the ejected material will reach the substrate but this can be high (approaching 100% in some systems) so even this geometry provides high deposition rates. To provide controlled energy-assisted deposition for fundamental studies of the process, ion beam systems were developed in the 1980s. As mentioned earlier, the first of these was the apparatus devised by Weissmantel and his colleagues (Weissmantel 1976, Weissmantel et al. 1979) but many other groups produced similar systems (Harper and Cuomo 1982, Harper et al. 1982a, b, 1983, Sainty et al. 1984, Martin et al. 1983, 1984, Colligon 1995, Kheyrandish et al. 1994). Figure 9.6 shows some of the various possibilities for ion-assisted deposition and includes an e-beam system for evaporation plus partial ionization (PIB, Fig. 9.6(a)) (Bai 1990)), ion-beam deposition (IBD, Fig. 9.6(b)) (Colligon et al. 1976, Lawson et al. 1975, Armour et al. 1986, Sharples et al. 1987), dual ion-beam deposition (DIB, Fig. 9.6(c)) (Harper and Cuomo 1982, Harper et al. 1982a, 1982b, 1983, Sainty et al. 1984, Martin et al. 1983, 1984, Colligon 1995, Kheyrandish et al. 1994), single-beam plus evaporation (SBE, Fig. 9.6(d)) (Weissmantel 1976, Weissmantel et al. 1979), ion-beam-induced chemical deposition (IBC, Fig. 9.6(e)) (Harriott and Vasile 1988) and ionized cluster-beam deposition (ICB, see Fig. 9.1) (Gspann 1993, Takagi et al. 1977, Yamada and Takagi 1995, Yamada 1994). In many of the ion-beam systems the Kaufman source was used (Kaufman et al. 1982) which produces broad beam gaseous ion species with a relatively low energy spread over an energy range 50-2000 eV and with intensity of order 1-5 mA/cm 2. Most of the above techniques are well known and the principle of operation is clear from the diagrams. In IBD the ions are usually extracted from an ion-implanter source at energies of 10-20 keV and, after magnetic separation, are slowed down to about 10 eV using carefully designed ion-optical lenses to prevent beam divergence and to steer the ions
234
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away from the direct line of sight where fast neutrals, created in charge-exchange collisions, would bombard and probably remove the deposited species (Colligon et al. 1976, Lawson et al. 1975, Armour et al. 1986, Sharples 1987). Such systems allow epitaxial growth of materials and can even grow isotopically pure films (Colligon et al. 1976, Lawson et al. 1975). In IBC a suitable organometallic vapor has to be directed at the substrate surface and the ions simply disrupt this to give a deposit plus a volatile species. Harriott
Physical Vapor Deposition
235
Table 9.2 Principal ion-beam-assisteddeposition methods. Method
Main features
Main application
PIB
Evaporated beam containing about 1% ionized species Direct deposition of 1-10 eV ions Few cm2, low rate One beam sputters One beam ion-assists Several cm2, low rate Precise control of IAD Evaporation plus ion-bombardment Several cm2, low rate Ion-beam dissociates chemical Vapor species at substrate Several cm2 Clusters of 100-1000 atoms About 1% are ionized Arrival energy about 10 eV/atom Several cm2
Epitaxial layers
IBD DIB
SBE IBC
ICB
Epitaxial layers isotopic films Novel compounds Metastable films
Novel compounds Metastable films Metal contacts Hard coatings Mask repair Metal contacts Epitaxial layers
and Vasile (1988) have used a fine-focused liquid-metal ion source with spot size just below 1 # m dia to accurately lay down carbon films and hence repair minute sections of damaged photolithographic masks. ICB systems operate by forcing an evaporant through fine nozzle which produces clusters of atoms of 2-5 nm size containing some 500-2000 atoms. Atoms in such clusters appear to have a binding energy lower than that in the bulk material (Anderson 1976) and the clusters have a low melting temperature (Buttet and Borel 1976). An e-beam crosses the path of the clusters and causes ionization of one or more atoms in the cluster which is then accelerated by applying a negative voltage of about 1 keV to the substrate. On impact the energetic atoms in the cluster disperse across the substrate surface and can produce epitaxial growth even on substrates where there is a large lattice misfit with the film material (Lawson et al. 1975). Later improvements in the temperature control of the nozzle have allowed higher deposition rates of 10-100 nrn/s to be attained for Zn (Gspann 1993), Ag (Gspann 1993) and Mg (Hagena et al. 1994). A summary of these ion-assisted methods and their main application is given in Table 9.2.
9.3.5 Magnetron sputtering All the above ion beam processes are important for fundamental studies of the energyassisted coating process and can be used for small-area coating where film thickness is well below 1/zm. A much quicker coating technique is based on the magnetron sputtering system which has now been adapted to provide a degree of ion bombardment of the growing film. A series of magnetron systems are illustrated in Fig. 9.7. In the conventional dc
236
J. S. Colligon
magnetron, shown schematically in Fig. 9.7(a), electrons are constrained to follow spiral orbits about the lines of magnetic flux which follow paths parallel to and above the target surface. As these electrons pass through the region above the target they have collisions with gas atoms, some of which ionize. If the target is biased to a sufficiently negative voltage, some of these ions will be attracted to it and cause sputtering. Sputtered atoms pass back through the ionizing region and, again, some of these are also ionized. Others escape and will fall onto the substrate to form the new film. When insulators need to be sputtered an rf supply is used. However, the rate of deposition is lower for rf and an alternative technique has been described by Sproul et al. (1983), Yashar et al. (1997) in which a pulsed 70 kHz positive voltage is superimposed on the negative dc supply for approximately 4 #s of the 14.2/zs cycle. This technique prevents charging and arcing when sputtering insulating materials. Another variation of the magnetron sputtering technique was first described by Window and Savvides (Window and Savvides 1986a and b, Savvides and Window 1986) and is known as the unbalanced magnetron. Here the magnetic fields are arranged so that some flux lines flow out towards the substrate. This can be achieved by using magnets of different area or strength (Fig. 9.7(b)), or, by applying an additional electromagnetic field using current-carrying coils (see Fig. 9.7(c)). Some electrons will now spiral along these field lines towards the substrate. Ions can be formed along the electron path and, also, because there is a net movement of negative charge from the magnetron region, positive ions will also be attracted to the substrate. The substrate will thus be bombarded by energetic ions and electrons, depending on its potential. A better practical system is to arrange for two (or four, or more) magnetrons to face each other in pairs having opposite magnetic polarities (Fig. 9.7(d)) and place the substrate in between them. Such systems, often described as closed field unbalanced magnetron sputtering systems (CFUBMS) have been described by Musil et al. (1990), Kadlec et al. (1990) and Teer (1990), who showed that there is a high probability of ionization of atoms in such configurations and that the substrate will be bombarded by ions of both the gas and film material. Under typical operating conditions, with chamber pressures of order 0.2-1 Pa, substrate ion bombardment currents of up to 100 mA at energies 10-100 eV are attained in such systems. Magnetron systems can give reasonable control of the energy-assisting contribution of depositing species and are capable of high rate deposition of the order of microns per minute. They do not allow independent control of these parameters however since the operating pressure influences the discharge conditions and, at higher pressures, collisions of particles with gaseous species can occur which change the energy of particles falling onto the substrate. More recent investigations by Ivanov et al. (1994) show that there is also additional bombardment of the substrate by particles reflected from the magnetron cathode. In spite of this, these authors have also shown that the presence of an external coil, which provides an additional magnetic field, can give reasonable control of deposition conditions. Improved adhesion of unbalanced magnetron sputtered coatings has been obtained by Miinz et al. (1989, 1991) using a preliminary arc bond sputtering process (ABSTM). In this system the coating process starts with a magnetron-assisted arc discharge where either Ar ions or a metal (Cr) ion first bombard the substrate at 800-1200 eV to give a
Physical Vapor Deposition
237
subsfrafe(floating or-V) I/Jat
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anced magnetron, (c) Unbalanced magnetron with controlled field, (d) Closed field Unbalanced magnetron.
preliminary etch. The main deposition then continues in the CFUBMS mode producing a coating which, for Cr-nitride, withstands a critical load over twice the value for coatings produced without the arc etch stage (Hurkmans et al. 1997). When sputtering chemically active species a problem can arise as a result of target poisoning; this is known as the hysteresis effect. It is particularly important for sputtering of titanium to form titanium nitride films since the Ti itself combines with, and pumps,
238
J. S. Colligon
the nitrogen and the speed of this additional pumping depends on the degree of poisoning of the Ti target. The problem can be overcome by using a high-speed pump on the system, by pulsing the nitrogen gas supply, or by using an optical monitor tuned to a Ti emission line and a feedback loop to the gas flow control. In the latter case, any reduction in Ti can be counterbalanced and stoichiometric TiN will continue to form.
9.4. I N F L U E N C E OF E N E R G Y ON COATINGS
The additional energy and the sudden cooling of the atoms forming thin films during ion-assisted deposition has the following effects: (1) film morphology changes and film density increases, (2) residual gases are desorbed from the substrate surface, chemical bonds may be broken and defects created, and hence nucleation processes are modified and film stress and adhesion changes, (3) for films containing two or more elements, the film composition can be chosen to produce a range of novel metastable phases which have application as hard coatings and low friction films, and (4) film microstructure is altered.
9.4.1 Film morphology and density One of the earliest correlations between coating morphology and substrate temperature (or adatom diffusion energy) was made by Movchan and Demchishin (1969). This model was developed further by Thornton (1974, 1977) to include the effect of chamber pressure. In the usual deposition situation, the majority of the condensing atoms arrive from one direction at the substrate. Depending on their own energy and the substrate temperature (T) they will have an opportunity to diffuse across the substrate surface to find a nucleation site. Subsequent atoms may join the earlier ones, or, if their energy is low or the temperature is low, may start new nucleation sites. At low energy and temperature, inverted pyramidal structures will grow toward the source of condensing atoms with large voids between them, forming the so-called Zone I structure. This occurs when the ratio of substrate temperature to the melting point of the condensate (Tm) is between 0.1 and 0.3. As the substrate temperature rises some additional surface diffusion can occur and the coating becomes more dense. For higher T~ Tm the growth process becomes dominated by the adatom surface diffusion and columnar crystalline grains separated by dense boundaries form. At even higher T~ Tm values bulk diffusion can also play a part leading to the formation of equiaxed grains. The additional energy provided to the condensing atoms during ion-assisted deposition can produce similar growth zone regimes but at lower substrate temperature since this transferred energy provides the adatoms with the ability to move further across the surface and to adjacent sites in the bulk. This is immediately obvious from the increase in film densities reported; for example Martin et al. (1984) found an increase in ZrO2 density of 14% when additional 600 eV oxygen ion bombardment was used compared to the density
Physical Vapor Deposition
239
for films formed by evaporation of Zr without bombardment. Computer simulations of Mtiller (1989) have predicted the increase in density for this same system extremely well.
9.4.2 Nucleation and adhesion
A review of ion-induced adhesion effects by the present author shows that there are many examples of improved adhesion resulting from energetic ion irradiation of films (Colligon and Kheyrandish 1989). There is no clear evidence however as to the prime cause for this improvement which could be a result of atomic mixing across the interface, or enhanced surface mobility leading to better bonding. Certainly, cross-sectional microscopy studies of films show that additional energy provided by ion bombardment can reduce and, even, eliminate columnar-type growth as shown by Petrov et al. (1989). This group varied the bias step by step during deposition in a dc planar magnetron system when sputtering TiN onto HSS plates. A substrate bias greater than 120 V was sufficient to produce fully dense films for substrate temperatures of 300~ but a bias exceeding 240-280 V was required and substrate temperature of 900~ before columnar growth was interrupted. It thus appears that it is a combination of substrate temperature and added energy, both leading to increased surface mobility of atoms, which will influence the film nucleation and subsequent morphology. If the ion/atom ratio is rather high then, as Rossnagel et al. (1982) and Robinson and Rossnagel (1984) have shown, not only do the atoms migrate on the surface to form island clusters, but also these clusters are more stable to re-sputtering so that conical features can form. At the nucleation stage, both Choi et al. (1990) and Tsai and Atwater (1992) have demonstrated improved growth resulting from a combination of increased surface mobility and dissociation of islands during the initial stages of deposition of semiconductor materials. This trend supports earlier work of Marinov (1977) who found that energetic condensation produced larger clusters of atoms owing to the break-up of smaller clusters and the increased surface mobility of the re-liberated atoms. As films become thicker, however, volume recovery processes start to dominate over surface energy, as seen by Ziemann and Kay (1983). They show that fcc metal films can have (111) faces parallel to the surface initially but, as a result of strain energy being reduced to a minimum, the orientation will change towards (100). In studies of the growth of epitaxial TiN films on MgO (100) substrates, Hultman et al. (1988) conclude that the following factors are significant: (1) at low temperatures point defects are not annealed out and may form dislocation loops, (2) for low energy ion-assistance this defect density is reduced as a result of enhanced atom mobility; and (3) for higher ion energy, defects are created by the energetic particles at a higher rate than they are annihilated and remain in the film. The competition between the various processes leads to an interplay between substrate temperature and ion energy, with the added proviso, noted by Sundgren et al. (1992), that the ion energy should only be sufficient to cause atom movement without introducing defects or gas bubbles. The use of higher energies for ion-assistance can lead to the im-
240
J. S. Colligon
plantation of ions which may introduce additional stress, form bubbles, and, according to Ziemann and Kay (1982), migrate to grain boundaries and influence grain size.
9.4.3 Metastable phases Some of the earliest experiments on the control of stoichiometry of ion-assisted films were reported by Hentzell et al. (1985). Using a dual ion-beam system, A1 atoms sputtered by one ion-beam were deposited onto a substrate which was simultaneously bombarded by 100-500 eV nitrogen ions. Depending on the arrival ratios of nitrogen to aluminium species the resulting A1Nx compound was formed. The quantity x was seen to increase with N/A1 ratio up to unity; for higher ratios the compound was always A1N. Similar studies reported by Rossnagel and Cuomo (1988) for 100 eV oxygen bombardment of a growing film of copper (in this case evaporated) showed that, as the O/Cu ratio increased, apparent stationary composition regions, corresponding to Cu20, Cu205 and CuO were produced. Further X-ray analysis indicated that the Cu20 and CuO samples were normal polycrystalline phases and, apparently, the Cu205 was a new metastable form of copper oxide, hitherto unseen. The appearance of a similar metastable form of TiNx has been seen by Kim (1991) using a dual-beam system to sputter Ti onto a sample simultaneously bombarded by N2 ions at 100-400 eV.
9.4.4 Microstructure With the potential control in nucleation, growth and atom mobility provided by the additional energy available in ion-assisted deposition processes it is not too surprising that some control of crystal structure and grain size in the films is now possible. Hultman et al. (1988) have shown that the defect structure in TiN grown epitaxially onto MgO(100) substrates is a function of the bombarding energy and target temperature. This defect level decreases as the ion energy increases in the region 100-200 eV but then increases again at higher energies as a result of the increased damage rate. A similar trend, but for fixed substrate temperatures (about 50~ was reported by Kheyrandish et al. (1994). In this work the least damaged crystalline grains, judged from the TEM patterns, were obtained for 400 eV N + ion bombardment energies (i.e. also 200 eV per N atom). Roy et al. (1997) have proposed that the surface packing density of a material dictates the crystal phase that will develop during energetic condensation; for Ni deposition they found that it was the (220) texture that becomes dominant during ion-assisted growth. Chadwick and Miller (1967) showed that the surface packing density is 0.8 for a random Ni hard-sphere liquid model, 0.556 for the (110), 0.907 for (111) and 0.785 for (100) planes. Thus, assuming the deposition process normally gives the random dispersion, some energy is needed to displace atoms to the (110) distribution. There is thus a correlation between added energy and grain size which may be exploited to tailor the properties of the coating, since there is also a correlation between grain size and hardness from the Hall-Petch relationship (Hall 1951, Petch 1953). In general, for grains of size above 100 nm, this relationship shows that alloys with smaller grain size exhibit higher strength. The added energy per atom, discussed above, has been shown by Harper et al. (1985) to be a useful guide for the prediction of deposition conditions leading to improved film
Physical Vapor Deposition
241
properties. Harper et al. plotted the results of a large selection of data from various sources on a diagram showing ion/atom arrival ratio against ion energy and found that most data came within the interval 1-100 eV/atom. A similar diagram has been presented by Martinu et al. (1994) but with data for SiO2, other dielectrics, metals and semiconductors identified separately. These average values, calculated from all atoms and all ions, although a good guide, do not always give a good prediction of the experimental result. This is shown by Petrov et al. (1993) for deposition of Ti0.5A10.5N alloys in a UHV unbalanced magnetron system. Figure 9.8, based on Petrov's work, shows that preferred orientation, composition and lattice parameter for the films was quite different when the ion-assist energy changed, even if the energy per atom was the same. This reinforces the trends already discussed by Ziemann and Kay (1982) and Roy et al. (1988). Roy et al. (1988) measured the effect on stress produced in 5 # m copper films as a function of added energy using 62 eV, or, 600 eV Ar + ions and found a marked difference in the effects for the same added energy when using high and low energy ions. The major problem with the concept of average energy per atom is that (a) the condensing particle energy is not all delivered to the top plane of atoms, and (b) even if it were, not all the atoms in the plane share this energy. Clearly there is an urgent need for molecular dynamics simulations of the process but, unfortunately, for high ion energies, this requires lengthy computer time. Mtiller's molecular dynamics calculation, mentioned earlier for ZrO2, was also applied to Ni films during ion-assisted deposition (MUller 1986a, 1986b, 1987) and shows, for 50 eV Ar + ions at 30 ~ to the normal and an ion/atom ratio of 0.16 on a growing Ni film, that an almost fully dense film is produced. Attempts to obtain a better average energy (but not necessarily at the surface) have been made by Brighton and Hubler (1987), who considered that, to produce dense films, each component film atom must be involved in a collision cascade. Hirsch and Varga (1978) attempted a similar calculation based on movement of atoms in a region influenced by a thermal spike created by the energetic particle, and Grigorov et al. (1988) proposed that, for any ion-assisted deposition situation, there is an optimum flux of ions required to produce best film quality and this flux corresponds to that required to produce one displacement per condensing atom. The situation is even more complicated by the fact that the energy deposited at the surface of the film has been shown by Berg and Nender (1990) and Berg and Katardiev (private communication) to depend strongly on the mass and density of the substrate material. This effect is seen to be particularly important at the start of film growth (or, during the start of growth of each component in a multilayer). The ultimate model must therefore take into account, for deposition at a given substrate temperature: (1) the energy deposited at the uppermost (and, perhaps, second) layer of the growing film; (2) the effect of lower layers (which may have different mass and density) on this deposited energy; and
J. S. Colligon
242
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0.417 -
/
0.i,16 0.4~5
0
:
I 50
I
I 100
(Ed) (eV/mefa[ afom}
Figure 9.8. Effect of added energy on Ti-A1-N films (from Petrov et al. 1993).
(3) the number of atoms receiving an energy above some threshold level. This could be a threshold for breaking bonds to allow surface diffusion, or, for stimulation of some chemical reaction. In the meanwhile, the best practical approach is to aim to use a high ion/atom ratio and condensing ion and atom energies as low as possible so that defects are not created
Physical Vapor Deposition
243
deep in the film. Fortuitously this is exactly the conditions experienced at the substrate in most unbalanced magnetron systems.
9.5. A P P L I C A T I O N S OF P V D C O A T I N G S
PVD coatings are used for a wide range of applications which include modification of the following properties: Optical: for laser optics, compact discs, anti-reflective coatings (for spectacles), optically absorbing coatings, mirrors, architectural glazing, selective solar absorbers. Chemical: corrosion-resistant coatings, catalytic coatings, sensor electrodes. Mechanical: hard coatings for cutting tools, low friction coatings, wear-resistant coatings. Electrical: electrical contacts and interconnects, solid-state devices, solar cells. Decorative: to provide attractive finish on many products.
Looking at the coating industry as a whole it is seen that PVD is not the main coating process and, when all coating methods are considered such as painting and electroplating, it is estimated that PVD represents about 4% of the total coating market (Gorham Newsletter). However this is 4% of a market in excess of 100 billion US$ per annum and still represents a large turnover. In addition, the ion-assisted processes are able to produce dense, well-bonded films and the process can run at relatively low substrate temperatures so that it allows coatings to be applied to materials with low melting points such as plastics. Perhaps the most important advantage of energy-assisted PVD is the ability to create novel phases of binary and ternary compounds and to produce materials with novel microstructural (and nanocrystalline) features. These unique film properties arise as a result of the rapid "cooling" of the fast condensing atoms forming the film which can approach 1014 K/s. Some typical examples of these special features in coatings with novel optical, chemical and mechanical properties are now described. Coatings applied for their electrical and electronic properties are normally specific to a certain small-scale requirement, such as mask repair, contacts, insulating layers and barrier layers and ion-beam methods are often employed. Coatings for decorative purposes have all the same requirements of durability and corrosion resistance as those described below but, because very large areas are to be treated (many square metres for glass and plastic roll-coating systems) a compromise has to be reached between the high throughput requirement and the controlled conditions described below for optimum quality of the coated product. In these systems ion-plating, arc ion-plating and magnetron sputtering are normally used.
9.5.10pticalcoatings
Methods for producing optical coatings have been reviewed by several authors (Martin and Netterfield 1986, Colligon 1992, Martin 1986, Gibson 1987) and it is clear that coating stability and performance is enhanced if they are fully dense and defect-free. Optimum conditions for forming some of the main optical materials, compiled from data in references (Allen 1983, McNeil et al. 1984a, Martin et al. 1983, McNeil et al. 1984b, McNally
244
J. S. C o l l i g o n
Table 9.3 Energy-assisted thin film optical coating conditions (measurement at 550 nm except where indicated). Refraction indix SiO2
1.46
TiO2
2.5
ZrO2
2.05
Ta205
2.25 (350 nm) 1.68 (350 nm) 1.8 (351 nm) 2.4 2.11 (350 nm) 1.39
A1203 Y203 CeO2 HfO2 MgF2
Ion
Energy(eV)
Substrate temperature(~
Ref.
O 30 Ar (oxygen background) O 300
Various
Ar 600 (oxygen background) O 500 (oxygen background) O 300-1000
20
McNeil et al. (1984) McNeil et al. (1985) Martin et al. (1983) Allen et al. (1983), McNeil et al. (1984), McNeil et al. (1985), Martin et al. (1983)
275
McNally et al. (1985)
275
McNally et al. (1985)
< 80
Cole et al. (1984)
20 275
Netterfield et al. (1985) McNally et al. (1985)
O 300 (oxygen background) O 1200 O 500 H2O
700
200
20
Martin and Netterfield (1985)
e t al. 1985, Cole e t al. 1984, Netterfield e t al. 1985 and Martin and Netterfield 1985) are presented in Table 9.3. One of the problems with the energy-assisted optical coating is that, if the ion/atom ratio is too high, or the ion energy is high, there is significant re-sputtering of the film and this can selectively remove one of the components. In addition the ion species itself can become buried in the film. Figure 9.9 illustrates this problem for CeO2 films produced using 1.2 keV oxygen ion-assistance. As the oxygen ion current density is increased the films have a higher refractive index and are more stable to exposure to air. However, beyond a critical current density value, the refractive index falls again owing to either oxygen incorporation, or, preferential removal of Ce. Proper selection of process parameters is thus critical for optical films. The ion-assist technique not only produces superior quality optical coatings, but also, presents a new opportunity to "tailor" the films by sputtering two or more materials at the same time. Figure 9.10 shows data by Cevro (1995) for a single layer anti-reflective coating for 1050 nm wavelength radiation. This was, in fact, produced by ion-beam sputtering of a Ta205/SiO2 mixed layer but, with the geometry used, some reflected ions also bombarded the growing films. Similar data has been reported by Laird and Belkind (1992) and Vretenar (1990) for TiO2/SiO2 films, by Denetale and Harker (1987) for TiO2/A1203 and by several groups (Pond e t al. 1989, Russak e t al. 1989, Gilmore e t al. 1987) for ZrO2/SiO2 films. Further
245
Physical Vapor Deposition 2.5
j
2.t~
I
AD CeOz
2.3 x
2.2
.c_ ~
2.1
r~
2.0
_> D._
~
1.9 1.8 I
films not stable
I
I
1
I
250 500 750 1000 1250 Oz ion current density (pA cm-2) ,
pgssible incorporation w
or oxygen or voua creation
Figure 9.9.
Variation in refractive index of CeO2 a a function of oxygen ion-assist energy. Films marked (x) are not stable on exposure to air. (From work of Netterfield et al. (1985)).
controlled studies of this type of new optical coating should lead to many new composite films with unique properties.
9.5.2 Corrosion protective coatings Another area where film density is important is that of corrosion protection. Ion-beam methods for producing such coatings have been reviewed by Wolf (1992) who notes that stress is an important factor in this type of application. Tensile stress can be reduced by increasing the ion/atom ratio in many cases, for example in Cr films using Ar ion assisted deposition (Roy et al. 1989a, b). Again the novelty of energy-assisted PVD is that new combinations of materials can be studied and their corrosion properties evaluated. Work reported by Hsieh et al. (1992), evaporating Nb and Cr with 250 eV Ar + ion-assistance at an ion/atom ratio of 0.68, has shown that Nb-50%Cr films are the most dense and column free but, because of its ability to form a Nb-rich oxide, Nb-30%Cr films had the best corrosion resistance. The present author has produced a range of Fe-Cr alloys for experiments which showed that the Cr concentration had to exceed 16% before complete passivation occurred (Fujimoto et al. 1993). The ability to "seed-in" additional elements has been exploited by Donahue et al. (1997) who showed that the addition of yttrium to magnetron-sputtered films of Ti-A1Cr-N drastically reduced the oxidation rate at elevated temperature.
246
J. S. Colligon
-
a~ ./Taz
0 s film
2.2
2.11 "0 e-
>
....
1.8
t.l
tU
"
1.6
1.~
x,,,~.,.,~,...~~ Si 0 z f i l m I , , , 250
Ta 1~%1 Si196%] I
I I 500
I
I
wavelength
I
I I 750
i
,
,
1 1000
-
(nm)
Figure 9.10. Refractiveindex of mixed SiO2 "Ta205 filmsas a function of wavelength (from work of Cevro (1995)).
9.5.3 Hard coatings Some of the main coatings used to produce durable hard layers include the nitrides of Ti, Zr, Hf and Cr, more complex nitrides which include Ti-A1-N, Ti-Zr-N, cubic BN, Ti-Nb-N, Ti-A1-O-N, diamond and diamond-like carbon, carbides of Ti, W and Ta and oxides. Hard coatings based on carbon, often referred to as diamond-like carbon (DLC) coatings, have been investigated using ion-assist methods since the early work of Aisenberg and Chabot (1971) who noted that diamond-like films were formed by directing a 40 eV beam of positive carbon ions at a substrate. Weissmantel (1981) later showed that similar films could be formed using ion-assisted techniques. More recently He et al. (1994) have shown that hard films with a high sp 3 content can be formed by sputtering graphite with Ar + ions onto silicon or AISI 52100 bearing steel. The growing film is bombarded with ions formed by introducing CH4 into a second source and running at energies from 0.2 to 25 keV. The highest sp 3 content occurred at the lowest (200 eV) energy where a Knoop hardness of 5100 kgf/mm 2 was registered using a 20 g load. A similar energy requirement was seen by Yugo (1994) in a microwave plasma running in a CH4/H2 mixture but here a two-step process was required, first a high level (40%) of CH4 was used to develop a carbon cluster and then, after 30 minutes, the CH4 level was reduced to only 0.5% allowing production of sp 3 bonds. This extensive field of DLC coatings has been reviewed by Ensinger (1996) who, in other work (Ensinger and Wolf 1989), has shown that these films provide good corrosion protection when applied to alloy steels. Silicon-carbide coatings have also attracted con-
247
Physical Vapor Deposition
1.2 O
-
o m
to C O
0.8
x
C:Si
~ ~ . .
~, A r ' S ( j X
v
"G 0.~ 0 a. E 0
0
, . . . . .
-4 . . . . .
0
100
-r . . . .
200
~mbard~ng ion energy
't-300 (eV)
~00
Figure 9.11. Compositionof ion-assisted Si-C films as a function of CH4 ion energy (from work of He et al. (1995)).
siderable interest (Learn and Haq 1970, Seaward et al. 1986, Yamada 1989, Berezhinskii et al. 1990, Yen and Murray 1990, He et al. 1995). Figure 9.11 shows data by He et al (1995) for dual beam ion-assisted growth using CH4 gas in the ion-assist source and Ar + ions to sputter silicon onto the substrate. It can be seen that the ion-assist energy must exceed 300 eV before the SiC compound forms. At this energy there was considerable re-sputtering of material and hence a low net rate of film growth. Cubic boron-nitride (c-BN) is an extremely useful coating for industrial applications owing to its high hardness, chemical and thermal stability at extremely high temperatures and low reactivity with ferrous alloys. It also has a large band-gap (>6 eV) and can be doped p-, or, n-type. Unfortunately, to form c-BN using conventional methods, pressures of over 11.5 GPa and temperatures in excess of 2000 K are required, which would destroy the properties of most substrates. However, ion-assist methods allow these films to be formed in vacuum at 460~ (see, for example Park et al. (1997)). In all cases a mixture of argon and nitrogen ions is required with an energy of about 300 eV and an ion/atom ratio of about 1.5. According to a model of the process by McKenzie (1993) and simulation experiments by Dworschak et al. (1995) the compressive stress, induced by the ions colliding with the film atoms, and the formation of interstitials is assumed to trigger the conversion of hexagonal-BN to c-BN. One of the most interesting aspects of current PVD methods is their ability to produce coatings with grain sizes smaller than those previously possible. This offers the possibility of producing nanoscale composite metastable alloys and pure materials which have unique properties including high hardness and high ductility. When the dimensions of components of the film become less than 100 nm the term/3 in the Hall-Petch equation
H = Ho + fld -1/2
248
J. S. Colligon
where H is the hardness, H0 a material constant and d the grain size, no longer applies since, according to Carsley et al. (1995) the material now has a significant fraction of its atoms in a grain boundary phase and so must be considered as a two-phase material. By considering typical grain boundary widths and grain sizes the fraction of each phase can be estimated and the overall flow stress of the composite material calculated, using bulk values of flow stress for the grains and amorphous material values for the boundary regions. The predicted values for hardness were compared by Carsley et al. with experimental data of Hughes et al. (1986) and E1 Sherik et al. (1992) for nickel and show good agreement. Similar agreement was also seen for iron and also copper. The maximum hardness is predicted to occur for grain sizes of about 4 nm for these materials. Vep~ek et al. (1995, 1996), Vep~ek and Reiprich (1996) and Christiansen et al. (1998) have predicted that mixed phase nitride materials should produce hard coatings and have demonstrated significant increases in hardness for crystalline TiN in amorphous Si3N4 formed by plasma-assisted CVD. Maximum hardness of 50-70 GPa was found for grain sizes of order 6-7 nm. Other combinations of nitrides are predicted to exhibit similar high values of hardness and Musil and Vlcek (1998) have noted that there are two types of mixed-nitride coatings: nanocrystalline metal nitride/hard matrix (such as c-TiN/a-Si3N4) and nanocrystalline metal nitride/soft matrix (such as c-TiN/Ni, Cu, Co, Ag, Au, rare earth elements). The metal nitride can be of several forms, including TiN, A1N, ZrN, NbN and WN. The former types of mixed nitride coatings are very hard but are brittle whereas the latter are of similar hardness but are ductile (Musil and Regent 1998, Musil and Zeman 1998, Nakayama et al. 1994). Musil et al. have shown that these mixed nanocomposite coatings can be formed using magnetron sputtering if the substrate is heated sufficiently to allow the second, amorphous phase, to migrate to the grain boundaries of the first phase (Musil and Vlcek 1998, Nakayama et al. 1994). It is this segregation that is believed to be responsible for limiting the growth in grain size. The value of substrate temperature required for formation of nanocomposite ZrN/Cu was found by Musil and Vlcek (1998), Musil and Regent (1998) to be of order 600~ and it is hoped that this can be reduced by improving the technique, for example, by incorporating a small quantity of a third element such as yttrium, which is known to reduce the grain size in metals (Liu et al. 1997). Another method for producing hard coatings is to use multilayer films A-BA-B . . . . , where the Young's modulus of one material is lower than that of the other. Seitz and Koehler (1956) predicted that the hardness would also increase for these "superlattice" structures and this prediction has been confirmed by Helmersson et al. (1987), for TiN/VN multilayers, by Kitagawa et al. (1998) for TiN/Ti multilayers and by Wu et al. (1997) for CNx/ZrN multilayers. 9.6. F U T U R E
TRENDS
The energy-assisted deposition processes are thus now producing a range of novel coatings which are well-bonded and can increase the lifetime of the coated product by significant factors. Coatings are formed mainly by magnetron sputtering, arc-ion-plating or other plasma-based techniques and it is expected that these methods will further improve to provide more efficient operation (for example, better utilisation of targets in magnetron systems) and better control of deposition parameters (for example addition of
Physical Vapor D e p o s i t i o n
249
electromagnetic coils to control the degree of ionization in magnetron sputtering). However the main further development appears to be in the production of new metastable alloy coatings (Bates and Arnell 1994, Monaghan and Arnell 1991) and nanocomposite coatings (Hughes et al. 1986, E1-Sherik et al. 1992, Vep~ek et al. 1995, 1996, Vep~ek and Reiprich 1996, Christiansen 1998, Musil and Vlcek 1998, Musil and Regent 1998, Musil and Zeman 1998, Nakayama et al. 1994) According to Musil and Vlcek (1998), ions bombarding a surface during coating impart sufficient energy by momentum transfer to allow high temperature phases to form within the collision cascade. The surrounding atoms however provide fast cooling at rates of order 1014 K/s so that the high temperature phase is frozen into the solid. Thus the presence of (a) an additional small quantity of several additional elements in the film, (b) additional energy to provide the effective temperature and pressure conditions within the film to form the high temperature phase, and (c) high cooling rates, will allow many phases to be formed at relatively low temperatures. Ion-assisted coatings can provide all the above parameters and the immediate work needs to study what types of coating can be formed and how stable they are in working conditions. The possible control of composition, crystal structure and stress in the coatings will lead to some exciting new possibilities in the near future, perhaps even the formation of the theoretical hard beta-phase cubic C3N4, predicted by Liu and Cohen (1989), Cohen (1994) and Yu et al. (1994) and currently occupying the attention of many research groups (Sj6str6m et al. 1996, Vep~ek et al. 1995, Riviere et al. 1995, Boyd et al. 1995).
ACKNOWLEDGMENT
The author wishes to thank many colleagues for providing information for this review.
REFERENCES Ahmed, N. A. G. and Teer, D. G. (1984) J. Phys. E, Sci. Instrum., 17, 411. Aisenberg, S. and Chabot, R. (1971) J. Appl. Phys., 42, 2953. Allen, T. H. (1983) Proc.Int. Eng. Cong., p. 1305. Anderson, A. B. (1976) J. Chem. Phys., 64, 4046. Armour, D. G., Bailey, P. and Sharpies, G. (1986) Vacuum, 36, 769. Bai, E, Yang, G.-R., Lu, T.-H. and Lau, L. W. M. (1990) J. Vac. Sci. Technol., A8, 1465. Bates, R. I. and Amell, R. D. (1994) Surface and Coatings Technol., 68/69, 686. Berezhinskii, L. I., Vlasnika, S. I., Lyashenko, G. I. and Rodionov, V. E. (1990) Phys. Stat. Solidi, (a)121, K193. Berg, S. and Katardiev, I. V. private communication. Berg, S. and Nender, C. (1990) in Handbook of Plasma Processing Technology. eds. Rossnagel, S. M., Cuomo, J. J. and Westwood, W. D. (Noyes, Park Ridge), p. 409. Berghaus, B. (1938) UK Patent 510, 993. Boyd, K. J., Martin, D., Todorov, S. S., A1-Bayati, A. H., Kulik, J., Zuhr, R. A. and Rabalais, J. W. (1995) J. Vac. Sci. Technol., 13, 2110. Brighton, D. R. and Hubler, G. K. (1987) Nucl. Instrum. Methods, B28, 527.
250
J. S. C o l l i g o n
Broszeit, E., Kloos, K. H. and Roth, Th. (1985) in Proc. 5th Int.Conf on Ion and Plasma Assisted Techniques. (CEP Consultants, Edinburgh), p. 337. Bunshah, R. E (1982) ed., Deposition Technologies for Films and Coatings. (Noyes, Park Ridge). Buttet, J. and Borel, J-P. (1976) Phys. Rev., A18, 2287. Carsley, J. E., Ning, J., Milligan, W. W., Hackney, S. A. and Aifantis, E. C. (1995) Nanostructured Mater, 5, 441. Carter, G. and Colligon, J. S. (1968) Ion Bombardment of Solids. (Heinemann, Oxford). Cevro, M. (1995) Thin Solid Films, 258, 91. Chadwick, G. A. and Miller, W. A. (1967) Metal. Sci. J., 1, 132. Choi, C-H., Hultman, L. and Barnett, S. A. (1990) J. Vac. Sci. Technol., A8, 1587. Chopra, K. L. (1969) Thin Film Phenomena. (McGraw-Hill, New York). Chow, E E (1991) in Thin Film Processes II, eds. Vossen, J.L. and Kern, W., Ch 2 and Ch 3, 1991, Acad. Press Inc., San Diego. Christiansen, S., Albrecht, M., Stunk, H. P. and Veprek, S. (1998) J. Vac. Sci. Technol., B16, 19. Cohen, M. L. (1994) Solid State Commun., 92, 45. Cole, B. E., Moravec, T. J., Ahmen, R. G. and Ehlert, L. B. (1984) J. Vac. Sci. Technol., A2, 372. Colligon, J. S. (1992) Vuoto, 22(3), 3. Colligon, J. S. (1995) J. Vac. Sci. Technol., A13, 1649. Colligon, J. S. and Kheyrandish, H. (1989) Vacuum, 39, 705. Colligon, J. S., Grant, W. A., Williams, J. S. and Lawson, R. P. (1976) in Proc Conf. on Applications of Ion Beams to Materials. (Institute of Physics Publishing, London), Conference series 28, p. 357. Conrad, J. R., Dodd, R. A., Han, S., Madapurn, M., Scheuer, J., Sridharan, K. and Worzala, E J. (1990) J. Vac. Sci. Technol., A8, 3146 Denatale, J. E and Harker, A. B. (1987) MRS Symp. Proc., 77, 181. Donahue, L. A., Smith, I. J., Mtinz, W.-D., Petrov, I. and Greene, J. E. (1997) Surface and Coatings Technol., 94-95, 226. Dworschak, W., Jung, K. and Ehrarut, H. (1995) Thin Solid Films, 254, 63. E1-Sherik, A. M., Erb, U., Palumbo, G. and Aust, K. T. (1992) Scripta Metall. et Mater, 27, 1185. Enomoto, Y. and Matsubara, K. (1975) J. Vac. Sci. Technol., 12, 827. Ensinger, W. (1996) Surface and Coatings Technol., 84, 363. Ensinger, W. and Wolf, G. K. (1989) Mater Sci. E n g . , A l l 6 , 1. Erttirk, E. and Heuvel, H. J. (1989) in Plasma Surface Engineering. eds. Broszeit, E., Mtinz, W. D., Oechsner, H., Rie, K-T. and Wolf, G. K., 1st Conf. on Plasma Surface Engineering, 19-23 Sept. 1988, Informationsgesellschaft Verlag, p. 553. Faraday, M. (1857) Phil. Trans., 147, 145. Foxon, C. T. and Joyce, B. A. (1981) Fundamental Aspects of MBE, in Current Topics in Materials Science, vol. 7, ed. E. Kaldis, North Holland, Amsterdam. Fujimoto, S., Newman, R. C., Smith, G. C., Kaye, S. P., Kheyrandish, H. and Colligon, J. S. (1993) Corrosion Sci., 35, 51. Gibson, U. J. (1987) Physics of Thin Films, vol. 13, ed. Francombe, M. H. and Vossen, J. L. (Acad. Press Inc., New York), p. 109. Gilmore, C. M., Quinn, C., Quadri, S. B., Gossett, C. R. and Skelton, E. E (1987) J. Vac. Sci. Technol., A5, 285. Goodman, C. H. L. and Pessa, M. V. (1986) J. Appl. Phys., 60, R65. Gorham Newsletter, No 1, October 1995, p. 3. Grigorov, G. I., Martev, I. N., Langeron, J-P. and Vignes, J-L. (1988) Thin Solid Films, 161, 249. Grove, W. R. (1852) Phil. Trans. Roy. Soc. (Lond.), 142, 87. Gspann, J. (1993) Nucl. Instrum. Methods in Phys. Res., B80/81, 1336. Gtinther, K. G. (1958) Z. Naturforsch., 13A, 1081. Hagena, O. E, Knop, G., Fromknecht, R. and Linker, G. (1994) J. Vac. Sci. Technol., A12, 282. Hall, E. O. (1951) Proc. Phys. Soc., B64, 747. Harper, J. M. E. and Cuomo, J. J. (1982a) J. Vac. Sci. Technol., 21,737. Harper, J. M. E., Cuomo, J. J. and Kaufman, H. R. (1982b) J. Vac. Sci. Technol., 21, 73. Harper, J. M. E., Cuomo, J. J. and Hentzell, H. T. (1983) Appl. Phys. Lett., 43, 547. Harper, J. M. E., Cuomo, J. J., Gambino, R. J. and Kaufman, H. R. (1985) Nucl. lnstrum. Methods in Phys. Res., B7/8, 886. Harriott. L. R. and Vasile, M. J. (1988) J. Vac. Sci. Technol., B7, 1035.
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He, X., Li, W. and Li, H. (1994) Vacuum, 45, 977. He, Z., Inoue, S., Carter, G., Kheyrandish, H. and Colligon, J. S. (1995) Thin Solid Films, 260, 32. Helmersson, U., Todorov, S., Barnett, S. A., Sundgren, J-E., Markert, L. C. and Greene, J. E. (1987) J. Appl. Phys., 62, 481. Henderson, E. (1977) in Proc. Int. Conf. on Ion and Plasma Assisted Techniques. (CEP Consultants, Edinburgh), p. 122. Hentzell, H. T. G., Harper, J. M. E. and Cuomo, J. J. (1985) J. AppL Phys., 58, 556. Herman, M. A. and Sitter, H. (1989) Molecular Beam Epitaxy: Fundamentals and Current Status. (SpringerVerlag, Berlin). Hirsch, E. H. and Varga, I. K. (1978) Thin Solid Films, 52, 445. Holland, L. (1956) Vacuum Deposition of Thin Films. (John Wiley & Sons, Inc., New York). Hsieh, J. H., Wu, W., Erck, R. A., Fenske, G. R., Su, Y. Y. and Marek, M. (1992) Surface and Coatings Technol., 51,212. Hughes, G. D., Smith, S. D., Pande, C. S., Johnson, H. R. and Armstrong, R. W. (1986) Scripta Metall., 20, 93. Hultman, L., Barnett, S. A., Sundgren, J-E. and Greene, J. E. (1988) J. Crystal Growth, 92, 639. Hurkmans, T., Lewis, D. B., Brooks, J. S. and Miinz, W.-D. (1997) Surface and Coatings Technol., 86--87, 192. Ivanov, I., Kazansky, E, Hultman, L., Petrov, I. and Sundgren, J-E. (1994) J. Vac. Sci. Technol., A12, 314. Kadlec, S., Musil, J. and Munz, W. D. (1990) J. Vac. Sci. TechnoI., AS, 1318. Kaufman, H. R., Cuomo, J. J. and Harper, J. M. E. (1982) J. Vac. Sci. Technol., 21, 725. Kheyrandish, H., Colligon, J. S. and Kim, J-K. (1994) J. Vac. Sci. Technol., A12, 2723. Kim, J-K. (1991) Thesis: Synthesis and characterization of titanium nitride films using a dual ion beam technique, University of Salford, July 1991. Kitagawa, M., Sayo, A., Kuroda, Y., Kusano, E., Nanto, H. and Kinbara, A. (1998) J. Vac. Soc. Japan, 41, 115. Laird, R. and Belkind, A. (1992) J. Vac. Sci. Technol., 10, 1908. Lardon, M., Buhl, R., Signer, H., Pulker, H. K. and Moll, E. (1978) Thin Solid Films, 54, 317. Lawson, R. P., Freeman, J. H., Colligon, J. S., Grant, W. A., Nobes, M. J. and Williams, J. S. (1975) Nucl. Instrum.and Methods, 131, 567. Learn, A. J. and Haq, K. E. (1970) Appl. Phys. Lett., 17, 26. Liu, A. Y. and Cohen, M. L. (1989) Science, 245, 841. Liu, Y., Singh, R., Poole, K., Diefendorf, R. J., Harris, J. and Cannon, K. (1997) J. Vac. Sci. Technol., B15, 1990. Marinov, M. (1977) Thin Solid Films, 46, 267. Martin, E J. (1986) J. Mater. Sci., 21, 1. Martin, P. J. and Netterfield, R. E (1985) Appl. Optics, 24, 1732. Martin. P. J. and Netterfield, R. E (1986) in Progress in Optics, ed. Wolf, E. (Elsevier Science Publishers BV, Amsterdam). Martin, E J., Clarke, G., Lanford, W. A. and Sie, S. H. (1983a) Appl. Phys. Lett., 43, 711. Martin, E J., MacLeod, H. A., Netterfield, R. E, Pacey, C. G. and Sainty, W. G. (1983b) Appl. Optics, 22, 178. Martin, P. J., Netterfield, R. P. and Sainty, W. G. (1984a) J. Appl. Phys., 55, 235. Martin, P. J., Netterfield, R. E, Sainty, W.G. and Pacey, G. C. (1984b) J. Vac. Sci. Technol., A2, 341. Martinu, L., Klemberg-Sapieha, J. E., Kuttel, O. M., Raveh, A. and Wertheimer, M. R. (1994) J. Vac. Sci. Technol., A12, 1360. Mattox, D. M. (1963) J. Appl. Phys., 34, 2493. Mattox, D. M. (1965) Sandia Corp. Monograph SC-R-65-852. Mattox, D. M. and Kominiak, G. J. (1972) J. Vac. Sci. Technol., 9, 528. McKenzie, D. R. (1993) J. Vac. Sci. Technol., Bll, 1928. McLeod, E S. and Mah, G. (1974) J. Vac. Sci. Technol., 11, 119. McNally, J. J., A1-Jumaily, G. A. and McNeil, J. R. (1985) Opt. News, 131. McNeil, J. R., A1-Jumaily, G. A., Jungling, K. C. and Barron, A. C. (1984) Appl. Optics, 23, 552. McNeil, J. R., Barron, A. C., Wilson, S. R. and Hermann, W. C. (1985) Appl. Optics, 24, 486. Melngailis, J. (1987) J. Vac. Sci. Technol., B5, 469. Migauchi, E. and Hashimoto, H. (1986) J. Vac. Sci. Technol., A4, 933. Monaghan, D. E and Arnell, R. D. (1991) Surface and Coatings Technol., 49, 298. Movchan, B. A. and Demchishin, A. V. (1969) Phys. Met. Metallogr., 28, 83. Mtiller, K-H. (1986a) J. Appl. Phys., 59, 2803. Mtiller, K-H. (1986b) J. Vac. Sci. Technol., A4, 184. MUller, K-H. (1987) Phys. Rev., B35, 7906.
252
J. S. C o l l i g o n
Mtiller, K-H. (1989) in Handbook oflon Beam Processing Technology. eds. Cuomo, J. J., Rossnagel, S. R. and Kaufman, H. R. (Noyes, Park Ridge), p. 241. Mtinz, W. D., Schulze, D. and Hauzer, E (1991) Surface and Coatings Technol., 50, 53. Miinz, W.-D., Vannisselroy, K., Tietema, R., Hurkmans, T. and Keiren, G. (1989) Surface and Coatings Technol., 37, 111. Musil, J. and Kadlec, S. (1990) in Proc.7th.Confon Ion and Plasma-Assisted Techniques, (IPAT). (CEP Consultants Ltd., Edinburgh). Musil, J. and Regent, E (1998) J. Vac. Sci. Technol., in press. Musil, J. and Vlcek, J. (1998) Czech. J. Physics, 48(10), to be published. Musil, J. and Zeman, P. (1998) Vacuum, to be published. Nakayama, A., Yoshioka, T. and Nomura, T. (1994) in Surface Modification Technologies VII. eds. Sudarshan, T. S., Ishizaki, K., Takata, M. and Kanata, K. (The Institute of Materials, London), p. 315. Neave, J. H., Joyce, B. A., Dobson, P. J. and Norton, N. (1983) Appl. Phys., A31, 1. Netterfield, R. P., Sainty, W. G., Martin, P. J. and Sie, S. H. (1985) Appl. Optics, 24, 2267. Park, K. S., Lee, D. Y., Kim, K. J. and Moon, D. W. (1997) J. Vac. Sci. Technol., A15, 1041. Petch, N. J. (1953) J. Iron Steel Inst., 174, 25. Petrov, I. N., Hultman, L., Helmersson, U., Sundgren, J-E. and Greene, J. E. (1989) Thin Solid Films, 169, 299. Petrov, I. N., Adibi, E, Greene, J. E., Hultman, L. and Sundgren, J-E. (1993) AppL Phys. Lett., 63, 36. Ploog, K. and Graf, K. (1984) Molecular Epitaxy of lll-V Compounds: A Comprehensive Bibliography 19581983. (Springer, Berlin). Pond, B. J., DeBar, J. I., Carnigia, C. K. and Raj, T. (1989) Appl. Optics, 28, 2800. Ramalingam, S., Qi, C. B. and Kim, K. (1987) US Patent 4 673 477. Rigsbee, J. M., Scott, P. A., Knipe, R. K., Ju, C. P. and Hock, V. E (1986) Vacuum, 36, 71. Riviere, J-P., Textier, D., Delafond, J., Jaouen, M., Mathe, E. L. and Chaumont, J. (1995) Mater Lett., 22, 115. Robinson, R. S. and Rossnagel, S. M. (1984) in Ion Bombardment Modification of Surfaces. ed. Auciello, O. and Kelly, R. (Elsevier, New York), p. 299. Rossnagel, S. M. and Cuomo, J. J. (1988) Vacuum, 38, 73. Rossnagel, S. M., Robinson, R. S. and Kaufman, H. R. (1982) Surf. Sci., 123, 89. Roy, K. R. R., Trogolo, J. and Cuomo, J. J. (1997) J. Electronic Mater., 26, 1270. Roy, R. A., Cuomo, J. J. and Yee, D. S. (1988) J. Vac. Sci. Technol., A6, 1621. Roy, R. A., Yee, D. S. and Cuomo, J. J. (1989a) MRS Symp. Proc., 128, 23. Roy, R. A., Petzkie, R., Yee, D. S., Karasinski, J. and Boulding, A. (1989b) MRS Symp. Proc., 128, 17. Russak, M. A., Jahnes, C. V. and Katz, E. P. (1989) J. Vac. Sci. Technol., AT, 1248. Sainty, G., Netterfield, R. P. and Martin, P. J. (1984) Appl. Optics, 23, 1116. Savvides, N. and Window, B. (1986) J. Vac. Sci. Technol., A4, 504. Seaward, K. L., Barbee Jr., T. W. and Tiller, W. A. (1986) J. Vac. Sci. Technol., A4, 31. Seitz, E and Koehler, J. S. (1956) Solid State Phys., 3, 305. Sharpies, G., Bailey, P., Donnelly, S. E., Van den Berg, J. A. and Armour, D. G. (1987) in Proc. Int. Conf. on Ion and Plasma-Assisted Techniques. (CEP Consultants, Edinburgh), p. 31. Sigmund, P. (1969) Phys. Rev., 184, 383. Sigmund, P. (1971) Rev. Roum. Phys., 17, 969. Sj6str6m, H., Hultman, L., Sundgren, J-E., Hainsworth, S. V., Page, T. E and Theunissen, G. S. A. M. (1996) J. Vac. Sci. Technol., A14, 56. Sproul, W. D. (1983) Thin Solid Films, 107, 141. Sundgren, J-E., Hultman, L., Hakansson, G., Birch, J. and Petrov, I. (1992) Proc. MRS Fall Meeting, 268, 71. Takagi, T., Yamada, J., Yanagwa, K., Kunori, M. and Kobiyama, S. (1974) Japan J. Appl. Phys., suppl. 2, 1. Takagi, T., Yamada, I. and Sasaki, A. (1977) Thin Solid Films, 45, 569. Teer, D. G. (1983) in Coatings for High Temperature Applications, ed E. Lang, (Applied Science Publishers, London), p. 83. Teer, D. G. (1990) in Proc. 7th. Conf. on Ion and Plasma-Assisted Techniques, (IPAT). (CEP Consultants Ltd., Edinburgh), p. 145. Thornton, J. A. (1974) J. Vac. Sci. Technol., 11, 666. Thornton, J. A. (1977) Ann. Rev. Mater. Sci., 7, 239. Tsai, C. J. and Atwater, H. A. (1992) Proc. MRS Fall Meeting, 268, 127. Varasi, M., Misiano, C., Mancini, C. and Sarfari, P. (1985) in Proc. 5th Int. Conf. on Ion and Plasma Assisted Techniques. (CEP Consultants, Edinburgh), p. 451.
Physical Vapor Deposition
253
Vep~ek, S. and Reiprich, S. (1996) Thin Solid Films, 268, 64. Vep~ek, S., Haussmann, M. and Reiprich, S. (1996) J. Vac. Sci. Technol., A14, 46. Vep~ek, S., Reiprich, S. and Shizhi, L. (1995) Appl. Phys. Lett., 66, 2640. Vep~ek, S., Weidmann, J. and Glatz, E (1995) J. Vac. Sci. Technol., A13, 2914. Vossen, J. L. and Kem, W. (1978) eds., Thin Film Processes. (Academic Press Inc., San Diego). Vretenar, E (1990) Vacuum, 40, 1912. Vyskocil, J. and Musil, J. (1990) Surface and Coatings Technol., 43/44, 299. Weissmantel, C. (1976) Thin Solid Films, 32, 11. Weissmantel, C., Reisse, G., Erler, H. J., Henny, E, Beuvilogue, K., Ebersbach, U. and Schurer, C. (1979) Thin Solid Films, 63, 315. Weissmantel, Chr. (1981) J. Vac. Sci. Technol., 18, 179. Window, B. and Savvides, N. (1986a) J. Vac. Sci. Technol., A4, 196. Window, B. and Savvides, N. (1986b) J. Vac. Sci. Technol., A4, 453. Wolf, G. K. (1992) J. Vac. Sci. Technol., A10, 1757. Wu, M-L., Lin, X-W., Dravid, E V., Chung, Y-W., Wong, M-S. and Sproul, W. D. (1997) J. Vac. Sci. Technol., A15, 946. Yamada, H. (1989) J. Appl. Phys., 65, 2084. Yamada, I. (1994) in Proc. Symp. on Ion and Laser Processing for Advanced Materials (SILPAM-94). eds. Takagi, T. and Tomimasu, T. (Univ. of Tokyo Press, Tokyo), p. 55. Yamada, I. and Takagi, T. (1995) Handbook oflon Beam Processing Techniques. (Noyes, Park Ridge), p. 58. Yashar, E, Rechner, J., Wong, M. S., Sproul, W. D. and Barnett, S. A. (1997) Surface and Coatings Technol., 94-95, 333. Yen, M. Y. and Murray, P. T. (1990) J. Mater. Sci., 25, 4929. Yu, K. M., Cohen, M. L., Hailer, E. E., Hansen, W. L., Liu, A. Y. and Wu, I. C. (1994) Phys. Rev., B49, 5034. Yugo, S. (1994) J. Vac. Soc. Japan, 37, 569. Ziemann, P. and Kay, E. (1982) J. Vac. Sci. Technol., 21, 828. Ziemann, P. and Kay, E. (1983) J. Vac. Sci. Technol., A1, 512.
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Chapter 10 Chemical Vapor Deposition 10.1. Introduction 10.1.1 Presentation 10.2. Gas-Phase Transport and Reactivity 10.2.1 Non-Reactive Fluid Flow 10.2.1.1 Geometrical Effects 10.2.1.2 Influence of Operating Conditions 10.2.2 Reactive Flows 10.2.2.1 Thermodynamic Approach 10.2.2.2 Gas-Phase Mechanisms and Kinetics 10.2.2.3 Simulations of the Gas-Phase Composition 10.2.2.4 Nucleation and Growth of Solid Particles from the Gas Phase 10.3. Solid Phase Formation 10.3.1 Solid-Gas Thermodynamic Equilibrium Approach 10.3.2 The Structure of CVD Layers 10.3.3 The Driving Force for Crystal Growth from the Vapor Phase 10.3.4 Surface Mechanisms and Kinetics 10.3.4.1 Rate of Elementary Surface Processes 10.3.4.2 Estimation of the Kinetics of Surface Processes in CVD 10.3.5 Heterogeneous Nucleation 10.3.5.1 Nucleation Theories 10.3.5.2 Two- and Three-Dimensional Nucleation 10.3.5.3 Epitaxial Nucleation and Growth 10.3.5.4 Multi-component nucleation and chemical reactions 10.3.5.5 Nucleation Enhancement 10.3.5.6 Selective Vapor Deposition 10.3.5.7 Theoretical and Experimental Studies of Nucleation in CVD 10.3.6 Crystal Growth 10.3.6.1 Mechanisms of Crystal Growth 10.3.6.2 Classification of Crystal Faces 10.3.6.3 Crystalline Morphology 10.3.6.4 Morphology and Normal Growth Rate of Crystal Faces 10.3.6.5 Relation between Preferred Orientation and Morphology of Polycrystalline Films 10.3.6.6 Theoretical Studies of Crystal Growth from the Vapor Phase 10.4. Conclusions References
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Chapter 10 Chemical Vapor Deposition E TEYSSANDIER and A. DOLLET
10.1. I N T R O D U C T I O N
In vapor deposition, a solid material is deposited from gaseous molecules onto a substrate to form a coating. The process is nowadays classified into two basic families according to the initial proposition of Blocher (1967): physical vapor deposition (PVD) when coatings are prepared from pure condensation processes, and chemical vapor deposition (CVD) when deposits are formed by chemical reactions only. These two routes were developed separately and many modifications have since been introduced, so that this distinction is no longer strictly valid. A great variety of processes related to CVD now exist. Some were developed to lower the deposition temperature by introducing (i) a glow discharge (plasma-activated CVD, PACVD, or plasma-enhanced CVD, PECVD) or (ii) less thermally stable organometallic precursors (organometallic CVD, OMCVD). Some were developed on the basis of a specific devicewa laser beam can be used, either for gas phase excitation or substrate heating in the laser-enhanced CVD process (LECVD). Certain denominations refer to specific types of coatings and related processes, such as epitaxially grown films for electronic or optoelectronic applications by the metalorganic vapor phase epitaxy process (MOVPE), or infiltration of ceramic materials in fibrous preforms by the chemical vapor infiltration (CVI) process. Distinction is usually made between two basic types of reactors: hot-wall and coldwall reactors (Fig. 10.1). In cold-wall reactors, the parts to be coated are heated by various means: electromagnetic induction or joule effect for metallic parts; tungsten-halogen lamps (RTPCVD, rapid thermal processing CVD) for non-conductive materials. In the hot-wall case, the whole reactor, usually tube-shaped, is heated by means of a classical furnace. Though the same phenomena take place in both cases, they present characteristic features: deposition can occur at hot walls thus leading to gas phase depletion, and the hot part in a cold-wall reactor is responsible for large thermal gradients and significant buoyancy-driven convection. These points will be further detailed. A lot of industrial coatings are currently produced by the CVD process, in the field of electronic, optical, optoelectronic or mechanical applications, and a considerable body of literature as well as scientific journals are devoted to CVD. The reader can also refer to the non-exhaustive review articles already published on the subject (Bryant 1977, Hess et al. 1985, Jasinski et al. 1987, Jensen 1987, Besmann et al. 1989, Jensen 1989b, Wahl 1989, Kleijn 1991, Hampden-Smith and Kodas 1995). The present review will report on 257
F. Teyssandier and A. Dollet
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the thermally-activated process that excludes any kind of physical activation either in the gas phase or on the surface. The basic mechanisms involved in the CVD process will be reviewed and illustrated.
10.1.1 Presentation Difficulties in obtaining a comprehensive view of the growth process arise from the number of interdependent parameters such as temperature, pressure, composition of the initial gas phase, flow rate, type of reacting chamber, etc. and the complexity of the various steps. The steps of the individual mechanisms involved, which are summarized in Fig. 10.2, can be divided into two highly interdependent parts: (a) phenomena occurring in the gas phase, such as bulk transport of the species in a hydrodynamic flow with diffusion mechanisms due to concentration gradients and possibly homogeneous reactions or gravity effects; in a cold wall reactor one must furthermore take into account thermal diffusion (Soret effect), and (b) the surface phenomena including adsorption, diffusion, reaction and desorption. In this case a distinction has to be made between the nucleation period when interactions with the substrate take place, and the steady-state growth of the coating from the gas phase.
259
Chemical Vapor Deposition
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Figure 10.2. Basic mechanisms in the chemical vapor deposition process.
The range of variation of the different parameters is as follows. In the thermally activated process, the temperature domain usually ranges from about 100 up to 2000~ The lowest deposition temperatures can be reached only by the use of organometallic precursors, since inorganic precursors, which are much more thermally stable, require at least 800~ to be decomposed. The total pressure in the reactor extends from atmospheric pressure down to a few tenths of a Torr (1 Tort = 133 Pa), so, according to the large characteristic dimensions of the reactors a continuum description is generally appropriate. Furthermore, since the gas velocities are low (Reynolds number < 100) and the Mach number is small (Ma < 1) the flows are laminar and the gas is assumed to be incompressible. The wide variety of relevant scientific domains is one of the difficulties of the CVD process and the various above-mentioned phenomena will be detailed in this review. The presentation order corresponds to the occurrence of the phenomena, from the introduction of the reactive mixture at the reactor entrance to the growth on the surface to be coated.
10.2. G A S - P H A S E T R A N S P O R T AND R E A C T I V I T Y
In CVD, the various chemical elements required to grow the film are transported by molecules in a forced flow from the entrance of the reactor towards the hot surface. The multicomponent, diffusive and reactive fluid flow can be an extremely complex system, and both flow phenomena and homogeneous chemical reactions affect the film growth. The fundamental equations describing the transport properties in continuous media are well known (Bird et al. 1960) and they will not be reviewed here. For a comprehensive review of transport phenomena in CVD see Holstein (1992), Jensen (1994). We will first focus on the fluid flow, ignoring its potential chemical reactivity which will be subsequently discussed.
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10.2.1 Non-reactive fluid flow 10.2.1.1 Geometrical effects The reactor has a critical influence on the uniformity of the fluid flow. The direction of the flow with respect to the gravitational field as well as the reactor geometry induce specific flow patterns such as recirculation cells. Under mass-transfer-controlled conditions the local flux of active species and hence the thickness uniformity and structural homogeneity of the coating is highly dependent on the flow pattern. Transport phenomena were described for a long time by simple analytical models based on boundary-layer similarity and film theory. The validity of the boundary layer hypothesis is questionable in many CVD reactors (Rosenberger 1987) and the absence of such a layer has been demonstrated in several cases by 3D numerical modeling since its thickness eventually exceeds the dimensions of the reactor for typical operating conditions. A stagnant layer theory was also encountered in the CVD field. This non-physical theory, which allowed decoupling convective and diffusive phenomena in several CVD analytical models, originates from flow visualization experiments. Eversteyn et al. (1970) interpreted a TiO2 particle free layer adjacent to the susceptor as evidence for a stagnant layer. Laser Doppler velocimetry experiments (Talbot et al. 1980), as well as modeling (Fotiadis and Jensen 1990) have revealed that thermophoretic forces were responsible for driving the seed particles away from the substrate. None of the above-mentioned approaches gives a satisfactory description of the physical processes involved in CVD, but they have provided simple models which have successfully predicted qualitative trends. The same film theories have also been developed for heat transfer (thermal boundary-layer), and for mass transfer (diffusion boundarylayer) (Hitchman 1980). All these theories are now encompassed in the approach which solves the coupled partial differential equations by numerical integration using finitedifference, finite-volume or finite-element methods. 1D to 3D simulation models are required depending on the type of reactor encountered. Both vertical and horizontal reactors are used. In the absence of gravitational effects, flow patterns are equivalent. This is not the case in CVD reactors on Earth, where natural convection resulting from essentially thermally induced density gradients is superposed on forced convection. The influence of natural convection on the process may be quite different according to the disposition of the reactor orientation. Vertical axisymmetric reactors such as impinging jet, stagnation point flow or rotating disk reactors have been widely investigated on the laboratory scale. The ideal stagnation point and rotating disk flows have uniform vertical gas velocities producing a boundary layer of constant thickness in which the gases move radially outwards. Due to the axisymmetric reactor geometry, 1D simulation provides a reliable description of the flow phenomena which agrees very well with flow visualization experiments (Wahl 1977, Fotiadis et al. 1987, Wang et al. 1995). Forced and natural convection in vertical reactors are coaxial phenomena. They have the same direction in the upward forced-flow case which stabilizes the axial density gradient. Recirculations may however be present in this case because of radial temperature gradients. Downward non-isothermal forced-flow results in counter-current influences that usually generate strong recirculation cells (Wahl
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1977). These cells are often eliminated by increasing the inlet flow-rate thus improving the thickness uniformity. Flow in horizontal reactors is much more complex, and any horizontal density gradient in a gravitational field results in buoyancy-driven free convection without threshold (Rosenberger 1987, Gokoglu et al. 1989). Both longitudinal and transverse rolls are then generated and 3D modeling is in this case required (Ouazzani and Rosenberger 1990). Under mass-transfer-controlled conditions the thickness uniformity is highly dependent on the flow pattern. The deposition profiles calculated in Fig. 10.3 (Jensen 1989a, Moffat and Jensen 1986) for two different flow patterns show that growth rate is enhanced or reduced respectively in the domains facing downward or upward secondary flow directions. The same phenomena can take place at the edge of substrates, frequently leading to their rounding off. Mass-transfer inside a reactor depends on both the reactor enclosure and the operating conditions. The latter can be modified from one experiment to another, but the reactor geometry is selected once for all. The design of reactors can be optimized by modeling a priori the flow and heat transfer. These calculations are invaluable in obtaining a better understanding of their behavior for various operating conditions. Several basic options can thus be checked: the shape, distance between inlet and susceptor, inlet tube diameter, introduction of baffles. The influence of these reactor features has been extensively illustrated in the case of a cold-wall vertical reactor (Fotiadis and Jensen 1990).
10.2.1.2 Influence of operating conditions Once the reactor has been selected, the experimenter has to determine the most appropriate operating factors. Several parameters can be chosen and these include temperature, total pressure, nature and proportion of precursors, and total flow rate. The nature and shape of the parts to be coated are the imposed parameters. Our understanding of the basic phenomena involved in the CVD process provides us guidance for selecting or optimizing the operating parameters. Analytical models were first used for that purpose. The influence on the growth rate of the basic parameters such as total pressure or temperature can be easily schematized and therefore memorized thanks to fairly simple models, such as the one proposed by Van den Brekel (1977) which successfully predicts qualitative trends. To quantify the relative influence of gas phase diffusion resistance transfer against the surface reactions kinetic limitation, he introduced the so-called CVD number, which is in fact a simplification of the Damk6hler (surface) number when only first-order surface reactions are considered (Hess et al. 1985). The general behavior of a CVD reactor can also be deduced from dimensionless groups that arise from the scaling of the governing transport equations (Jensen 1989). These groups are also helpful in determining the influence of a specific parameter, such as temperature, on the physical properties of gases and thus the transport properties in CVD reactors (Rosenberger 1987). 10.2.2 Reactive flows In many situations, the residence time and gas temperature are high enough to allow reactions to occur in the gas phase prior to reaching the surface. The new species so
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Figure 10.3. Flow profiles and corresponding growth rate of GaAs calculated with a 3D simulation including a one species (TMGa) diffusion model with a fast surface reaction (presented in Jensen 1989a after Moffat and Jensen 1986). Left: insulation of side walls, right: cold side walls.
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produced are transported towards the substrate at variable rates according to their physical properties. Therefore, a large variety of species having different reactivities can be present in the vicinity of the deposition surface. The main CVD processing domains are traditionally reported in terms of growth rates and their dependence on temperature. Surface kinetics or gas phase mass-transfer control operating domains are thus deduced from the overall activation energy. To determine the apparent activation energy, it is usual practice to plot the logarithm of the growth rate against 1/T. Three temperature domains can thus be observed: (a) in the low-temperature domain, the deposition rate increases rapidly with temperature. The steep slope which is indicative of a high overall activation energy is associated with a growth rate governed by the surface phenomena. For very short residence times and/or very low temperature of the gas mixture, almost no intermediate species are formed, and the flow is assumed to be chemically frozen. (b) at intermediate temperatures, surface kinetics become fast and the ability of the process to supply reactive species to the surface becomes the growth-limiting factor. In this range of temperature, convection in the bulk and diffusion in the vicinity of the surface become the rate-limiting transport phenomena. Their weak dependence on temperature is responsible for the low overall activation energy. Higher gas phase temperatures allow homogeneous chemical reactions to develop under kinetic control. Near-equilibrium conditions are reached in the gas phase for long residence times and/or high temperatures. (c) at high temperatures, a decrease in the growth rate is often encountered. Several phenomena may account for such a behavior: the nucleation and growth of particles in the gas phase that deplete the amount of reactive species on the surface, modifications of the fluid flow dynamics or desorption of species that exceed the number of molecules impacting on the surface. In the intermediate and high temperature domains (cases (b) and (c)) gas phase phenomena and especially mass transport phenomena are the factors that control growth rate uniformity. The transition from the surface kinetics rate limiting domain to the mass transport limiting domain is of prime importance in the process as it leads to very specific features of the growth mechanisms. Gas-phase transport-limiting process is often responsible for surface roughening, whereas surface kinetics limitation tends to smooth out the surface (Van den Brekel 1978, Palmer and Gordon 1988, Palmer and Gordon 1989) (see also crystal growth section in part 3).
10.2.2.1 Thermodynamic approach Because the detailed reaction mechanisms and corresponding kinetic data are rarely available, chemical equilibrium has been the first traditional approach to CVD process modeling. Thermodynamic calculation has been applied to a great variety of chemical systems and for various purposes which will be detailed in the "Solid-gas thermodynamic equilibrium approach" section in part 3. Application of thermodynamic calculations to gas phase in the absence of solid state materials leads to pseudo-thermodynamic equilibrium involving a supersaturated gas phase. This allows determining bulk compositions and identifying key species, as well as conditions under which they might form (Allendorf 1993). Though this equilibrium situation, which is representative of a gas phase composition that would be reached at infinite residence time by a reactive mixture flowing at constant temperature (as long as homogeneous nucleation does not take place), is never
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F. Teyssandier a n d A. D o l l e t
encountered in CVD reactors, near- equilibrium states are often approached in hot-wall reactors. Accurate modeling and, ultimately, control of the CVD process requires knowledge of both the identity of the species formed during gas phase transport and their rates of formation. When these are known, the flux of reactive species toward the surface can be predicted. For that purpose, gas phase reaction mechanisms have to be identified and their corresponding rate constants must be measured or estimated.
10.2.2.2 Gas-phase mechanisms and kinetics In the dilute medium encountered in the CVD process, it is assumed that there are only two types of kinetic processes: unimolecular bond dissociation reactions and bimolecular reactions. For further details on experimental and theoretical development of gas-phase kinetics the reader can refer to classic texts (Benson 1976, Forst 1973, Gilbert and Smith 1990, Robinson and Holbrook 1996). Only a brief review of the basic principles will be given here. In a unimolecular reaction, an isolated molecule undergoes a chemical change. Reaching an internal energy sufficient for the reaction to occur requires excitation, which most frequently is obtained from collision with a bath gas at thermal equilibrium. Unimolecular reaction rates can be obtained experimentally from static (Neudorfl et al. 1980) and flow pyrolysis (Ring et al. 1970), shock tubes (Mick et al. 1996) and laser powered (Jasinski and Estes 1985) homogeneous pyrolysis techniques. They may also be estimated from theoretical models. The energy-dependent rate constant k ( E ) can be computed from the RRKM model (Rice-Ramsperger-Kassel-Marcus) which is the application of the transition state theory to a microcanonical ensemble of excited reactant molecules. This model gives an exact upper bound to the rate coefficient. A simpler model (QRRK) (Westmoreland et al. 1986) in which a molecule is treated as if all its vibrations can be represented by a simple average frequency (usually the geometric mean) can be used instead of the RRKM model. Three categories of bimolecular reactions (A + B --+ products) may be distinguished: displacement, association and metathesis reactions (in which an atom is transferred from A to B). The bimolecular adaptation of the unimolecular QRRK model (Dean 1985, Westmoreland et al. 1986) can be used to predict the pressure and temperature dependence of addition or recombination reactions that initially form excited intermediates. Once the various gas phase reaction mechanisms are identified and their rate measured or estimated, the kinetic data can be introduced in full-scale models or simpler (plug-flow or well-stirred) time-dependent models such as the Sandia CHEMKIN/SENKIN package (Kee et al. 1989, Lutz et al. 1995). These tools are useful in interpreting the evolution of gas mixtures (Teyssandier and Allendorf 1998).
10.2.2.3 Simulations of the gas-phase composition Even the simplest chemical systems may involve numerous gaseous species. The model proposed by Coltrin et al. (1989) for silicon deposition from silane, which includes only two elements (Si and H), is composed of 27 chemical reactions involving 17 gas phase chemical species. Adding carbon to allow SiC deposition, increases the gas phase chemical model up to 83 reactions involving 41 species (Allendorf and Kee 1991). Some of these species are not important in the deposition process and such systems can be simpli-
Chemical Vapor Deposition
265
fled for specific total pressure and temperature conditions (Wang et al. 1994). However the rejection of some species or chemical reactions according to concentration criteria can be misleading. As a fact, the selected subsystem must constitute a closed system with respect to the chemical reactions considered, i.e., the important reaction pathways for the selected species must not be ignored. Furthermore, due to the large difference of surface reactivity between closed-shell molecules and radicals, the latter species, even when present at low concentration, may have a significant influence in the surface process. For instance, though Sill2 forms only at high temperatures and is thus usually present at the interface with the surface substrate in far lower concentrations than the precursors introduced at the reactor entrance (silane or disilane), its unity reactive sticking coefficient (RSC) at the surface is at least three orders of magnitude larger at 1000 K than those of the saturated silanes (Sill4, Si2H6, Si3Hs, etc.). The contribution of silylene to silicon growth rate (at least at high temperatures) can be for that reason significant. Full-scale modeling including gas-phase reactions allows determining the contribution of the various gaseous species to the deposition rate. The partial deposition rates of Sicarrier species at the surface facing a circular impinging jet (Fig. 10.4) reveals that (i) the deposition rate increases rapidly towards the edge of the impingement surface as a result of the increasing amount of highly reactive intermediate species formed, (ii) the contribution of the initial precursor Si2H6 is constant on the surface that reflects its supersaturation which is a consequence of its low surface reactivity, and (iii) Si3H8 is responsible for the step variation of the deposit thickness observed at the edge of the surface.
10.2.2.4 Nucleation and growth of solid particles from the gas phase The non-controlled introduction of solid particles in a CVD reactor is detrimental to the properties of the films. These particles are usually introduced into the reactor as impurities or can be formed by agglomeration of smaller particles nucleated in the gas phase. They act as extrinsic seeds for heterogeneous nucleation, deplete the species necessary for the heterogeneous reactions and create undesirable microstructures including porosity that lower the mechanical properties of the deposits. The elementary processes involved in the formation and growth of particles from gaseous precursors include gas phase chemical reactions, homogeneous nucleation, heterogeneous condensation, coagulation and coalescence or fusion (Fig. 10.5). Both high temperatures and long residence times are favorable to homogeneous nucleation of such particles. This situation is especially encountered at the center of recirculation cells (infinite residence time) in hot-wall reactors. In contrast, the risk of solid particle formation is considerably reduced under low total pressure and/or in diluted reactive mixture. Because of the large temperature gradients found in cold-wall CVD reactors, the thermophoresis forces may cause the solid particles to deviate significantly from the fluid particle trajectories. This phenomenon is the preponderant mass-transfer mechanism used in the modified chemical vapor deposition process (MCVD) to manufacture optical waveguide preforms (Kim and Pratsinis 1988, Lin et al. 1990). Since the thermophoretic coefficient increases with decreasing pressure, its influence is expected to be enhanced with decreasing pressure (Fotiadis and Jensen 1990). Accordingly, low pressures, which reduce buoyancy-driven flows, favor rapid transport of particles away from the deposition zone and reduce particle contamination. The carrier gas has also a significant influence on
F. Teyssandier and A. Dollet
266
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Chemical Vapor Deposition PRODUCT
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267
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the thermophoretic transport, and under the same operating conditions the thermophoretic coefficient for nitrogen is smaller than that for hydrogen (Fotiadis and Jensen 1990).
10.3. S O L I D P H A S E F O R M A T I O N
Once introduced into the CVD reactor, the reactants may experience some chemical transformations during their transport towards the substrate. When reactants and reaction products reach the solid-gas interface, the gaseous chemical species undergo new complex physical and chemical interactions that may lead to film growth provided that some thermodynamic and kinetic criteria are satisfied. We will first discuss the energetics of the solid-gas transformation, and subsequently, the kinetics of the elementary heterogeneous processes. The nucleation and crystal growth phenomena will be finally addressed.
10.3.1 Solid-gas thermodynamic equilibrium approach The final equilibrium state of a thermally-activated CVD system, in which various chemical reactions are involved in the formation of a solid material from a gas phase, is imposed by thermodynamics whether or not this state is reached for kinetic reasons. In other words, when quantitative predictions are difficult to obtain because of the lack of mechanistic and kinetic information, thermodynamic calculations may provide reliable qualitative trends. Several review articles have already been devoted to this subject (Bernard and Madar 1990, Bernard and Madar 1991, Spear and Dirkx 1990). When all gaseous and solid species are taken into account, the "closed box" (nonflowing closed system) thermodynamic calculation at equilibrium provides valuable qualitative information on the nature of the different phases formed, on their amounts,
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as well as on the influence of the various parameters. Two methods have been
