Mechanical Alloying: Fundamentals and Applications

MECHANICAL ALLOYING

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MECHANICAL ALLOYING Fundamentals and Applications P.R. Soni Department of Metallurgical Engineering Malaviya Regional Engineering College, Jaipur, India

CAMBRIDGE INTERNATIONAL SCIENCE PUBLISHING iii

Published by

Cambridge International Science Publishing 7 Meadow Walk, Great Abington, Cambridge CB1 6AZ, UK http://www.cisp-publishing.com

First published 2001

© ©

P.R. Soni Cambridge International Science Publishing

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British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library

ISBN 1 898326568

Production Irina Stupak Printed by Pear Tree Press Ltd, Stevenage, England

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TO MY FATHER In memory of my father whom I loved most m He expired in February 1996

About the Author Dr. P.R. Soni received his M.Sc. (1975) degree in Physics from the University of Udaipur; M.E. (1979) and Ph.D. (1991) from the University of Rajasthan, Jaipur. Dr. Soni joined the Department of Metallurgical Engineering at Malaviya Regional Engineering College, Jaipur in 1980, where presently he is a Reader. He had been a Research Fellow at the Indian Institute of Technology, Bombay, in 1989. Dr. Soni was selected as Manager (R&D) of Nappro Synthetics in 1990. For the last fifteen years, he has been actively engaged in research and development activities related to mechanical alloying, composite materials and other P/M materials, and is closely associated with industries through consultancy. He has published thirty technical papers in various peer reviewed journals. He is a member of the APMI International and many other professional bodies.

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Contents 1 INTRODUCTION ......................................................................................... 1 1.1 HISTORY ........................................................................................................................... 1 1.2 BENEFITS OF MECHANICAL ALLOYING ............................................................... 3

2 MECHANICAL ALLOYING ........................................................................... 6 2.1 ALLOYING MILLS ......................................................................................................... 6 2.1.1 Mills in Practice .............................................................................................................. 6

2.1.1.1 Szegvari attritor mill....................................................................................... 6 2.1.1.2 Spex vibratory mill ......................................................................................... 8 2.1.1.3 Planetary ball mill .......................................................................................... 9 2.1.1.4 Large diameter ball mills ............................................................................... 9 2.1.1.5 Grinding media ............................................................................................. 10 2.1.2 Improved Mills .............................................................................................................. 10

2.1.2.1 Modified attritor ........................................................................................... 11 2.1.2.2. Uni-ball mill ................................................................................................ 12 2.2 THE PROCESS ............................................................................................................... 15 2.2.1 Process Monitoring ....................................................................................................... 19 2.3 FACTORS AFFECTING ................................................................................................ 19 2.3.1 Mill Parameters ............................................................................................................ 20

2.3.1.1. Impact energy .............................................................................................. 20 2.3.1.2 Size of the grinding ball ............................................................................... 20 2.3.1.3 Ball–to–powder ratio ................................................................................... 20 2.3.1.4 Speed ............................................................................................................ 20 2.3.2 Temperature .................................................................................................................. 21 2.3.3 Atmosphere ................................................................................................................... 22 2.3.4 Contamination ............................................................................................................... 23

3 VARIATIONS OF MECHANICAL ALLOYING .............................................. 25 3.1 REACTION MILLING .................................................................................................. 25 3.2 CRYOMILLING ............................................................................................................. 26 3.3 REPEATED ROLLING .................................................................................................. 26 3.4 DOUBLE MECHANICAL ALLOYING ...................................................................... 28 3.5 REPEATED POWDER FORGING .............................................................................. 29

4 PROCESS CONTROL AGENTS IN MECHANICAL ALLOYING ..................... 31 5 MECHANISMS IN MECHANICAL ALLOYING ............................................ 35 5.1 ALLOYING ..................................................................................................................... 35 5.1.1 Ductile–Ductile System ................................................................................................. 35 5.1.2 Ductile–Brittle System .................................................................................................. 36 5.1.3 Brittle–Brittle System ................................................................................................... 37 5.1.4 Idealness of MA Alloys ................................................................................................. 39 5.2 METASTABLE PHASE FORMATION ....................................................................... 40 5.2.1 Amorphization .............................................................................................................. 40 vii

5.2.2 Nanocrystallization ....................................................................................................... 49 5.2.3 Solid Solubility Extension (SSE) ................................................................................. 51 5.3 ACTIVATION OF SOLID STATE CHEMICAL REACTION .................................. 51

6 ENERGY TRANSFER AND ENERGY MAPS IN MECHANICAL ALLOYING ................................................................................................. 55 7 CONSOLIDATION OF MECHANICALLY ALLOYED POWDERS .................. 58 7.1 CONSOLIDATION TECHNIQUES ............................................................................. 58 7.2 THERMOMECHANICAL TREATMENTS ................................................................ 62

8 MECHANICAL PROPERTIES OF MECHANICALLY ALLOYED MATERIALS ................................................................................................. 65 8.1 TENSILE PROPERTIES ............................................................................................... 65 8.2 FRACTURE ..................................................................................................................... 69 8.3 CREEP ............................................................................................................................. 69 8.4 SCC SUSCEPTIBILITY ................................................................................................ 70

9 MODELLING MECHANICAL ALLOYING .................................................. 72 9.1 INTRODUCTION ........................................................................................................... 72 9.2 MECHANISTIC MODELS ........................................................................................... 72 9.2.1 Deformation, coalescence and Fracture ...................................................................... 73 9.2.2 Evolution of Particle Size ............................................................................................. 77 9.2.3 Milling Times ................................................................................................................ 79 9.2.4 Powder Heating ............................................................................................................ 85 9.2.5 Powder Cooling ............................................................................................................. 86 9.3 ATOMISTIC MODELS ................................................................................................. 87 9.4 THERMODYNAMIC AND KINETIC MODELS ...................................................... 89

10 JOINING OF MECHANICAL ALLOYING MATERIALS ............................... 91 10.1 WELDING ..................................................................................................................... 91 10.2 BRAZING ...................................................................................................................... 93 10.3 FORGE BONDING ....................................................................................................... 94

11 RAPID SOLIDIFICATION AND MECHANICAL ALLOYING ....................... 96 11.1 RAPID SOLIDIFICATION VERSUS MECHANICAL ALLOYING ..................... 96 11.2 MECHANICAL ALLOYING OF RAPIDLY SOLIDIFIED POWDERS ............... 96

12 APPLICATIONS ....................................................................................... 103 12.1 NICKEL-BASE SUPERALLOYS ............................................................................. 103 12.2 MA STEELS ................................................................................................................ 107 12.3 ALUMINIUM-BASE MATERIALS .......................................................................... 111 12.4 COPPER-BASE MATERIALS .................................................................................. 119 12.5 TITANIUM SYSTEM ................................................................................................. 121 12.6 MAGNESIUM-BASE MATERIALS ........................................................................ 123 12.7 SUPERSATURATED SOLUTIONS ......................................................................... 125 12.8 MAGNETIC MATERIALS ....................................................................................... 126 12.9 MA POWDERS FOR SPRAY-COATINGS ............................................................. 129 12.10 SUPERPLASTICITY ............................................................................................... 130 12.11 TRIBOLOGICAL MATERIALS ............................................................................ 130 12.12 COMPOSITES .......................................................................................................... 134 viii

12.13 AMORPHOUS SOLIDS ........................................................................................... 136 12.14 NANOCRYSTALLINE MATERIALS .................................................................... 137 12.15 MECHANICALLY ACTIVATED CHEMICAL REACTIONS ........................... 138 12.16 OTHERS .................................................................................................................... 140 CONCLUSIONS .................................................................................................................. 140 LIST OF SYMBOLS ........................................................................................................... 145 INDEX .................................................................................................................................. 147

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PREFACE From the total initial laboratory success in 1968, the process of mechanical alloying (MA) has been developed into a well-controlled production operation over the years, and applied to develop varieties of materials. Being a new field, there is a wealth of recent scientific literature available, but it is all scattered – The problem of a beginner to get started in a practical way with the MA technique. This book tries to address this problem and is aimed at the undergraduates, postgraduates, materials scientists and engineers who want to have in-depth knowledge in this field. The book is also designed to serve as an introductory and refreshment reference tool for the manufacturing engineers actively involved in MA or the allied industry but are in need of detailed knowledge of metallurgical engineering or materials science. A two year metallurgical engineering or materials science course should provide the necessary basis for comprehension of the material discussed in the book. This book tries to put forward the fundamentals of MA recipes where the technique has been successful and highlights the areas in technology where it can provide benefits in developing high-tech materials. Not only this, many secrets of the MA processing approach are still in the embryonic stage and this book creates a brainstorm in the mind of materials scientists and engineers to reveal the same. The book comprises of twelve chapters. Starting from the historical development of the MA technique and highlighting its benefits in Chapter 1, Chapter 2 discusses the basic process of MA devices used for, and factors affecting the process. This chapter deals with the different variations of MA which have been developed over the course of time. Chapter 4 deals with the process control agents used in MA, while Chapter 5 deals with mechanisms involved in basic processes of MA, metastable phases formation and activation of solid state chemical reactions. Chapter 6 deals with energy transfer and energy maps in MA. Chapters 7 and 8 deal with basics of consolidation of MA powders and mechanical properties of MA products, respectively. Chapter 9 explains the basics of models developed to predict the results of an MA process. Joining techniques for MA materials are discussed in Chapter 10. The two nonequilibrium processing techniques, MA and rapid solidification, are compared and combinations of the two to enhance the properties of the xi

Mechanical Alloying

rapidly solidified materials, have been discussed in Chapter 11. Chapter 12 is meant for highlighting the cases where the technique has been applied to produce MA products at industrial levels and the potential materials to find applications in the industrial scene. My sincere thanks to the Minerals, Metals and Materials Society, Warrendale, APMI (American Powder Metallurgy Institute) International, New Jersey; Dr. J.S. Benjamin, Hitchiner Manufacturing Company, Inc, Milford; Dr. H.K.D.H. Bhadeshia, University of Cambridge; Prof P. Ramakrishnan, Indian Institute of Technology, Bombay; Prof. F.H. Froes, University of Idaho and Prof. C.C. Koch, North Carolina State University, for their co-operation in completion of this task. I am grateful to Prof. T.V. Rajan, Malaviya Regional Engineering College, Jaipur for encouraging to pursue my interest in MA. I wish to acknowledge my wife Pramila and my children Anshu and Ankit, for bearing my disappearance when I was 'in the book'. P.R. Soni

J.S. Benjamin, 'Father' of mechanical alloying (1968).

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Introduction

1 INTRODUCTION 1.1 HISTORY The use of inert additions to improve elevated temperature mechanical properties of metals was first exploited in 1910 by W.D. Coolidge in thoriated tungsten [1]. The development of dispersion-strengthened alloys by internal oxidation started in 1930 [2] and the invention of dispersion-strengthened aluminium took place in 1949 [3]. However, the relatively low melting point of aluminium was a severe limitation for the use of SAP at elevated temperatures. This led to attempts in applying dispersion strengthening to higher melting point metals such as copper and nickel. In these metals, the self oxides cannot be used as they are not sufficiently stable against Ostwald ripening at elevated temperatures. ThO 2dispersed nickel, having a finely distributed dispersoid, was produced successfully by melting to improve the mechanical properties (1960) [4]. The material has vastly developed elevated temperature properties. However, the use of such materials was still limited due to their low strength at intermediate temperatures and lack of corrosion resistance. Hypothetically, these shortcomings could be solved by combining the corrosion resistance and intermediate temperature strength of γ′−precipitation-hardened, nickel-base superalloys with the high temperature strength and stability of oxide dispersion strengthening. One of the primary hurdles to materialize this idea is the production of uniform dispersion of fine oxide particles, less than 0.1 µm in size, in alloy powder particles in such a manner which leads to interparticle spacing of less than 0.5 µm in a consolidated product. The nickel-base superalloys contain chromium, aluminium and titanium for effective precipitation hardening, elements which are easily oxidizable. Aluminium and titanium oxides are so stable that they cannot be reduced to a metallic state, required in an alloy, without reducing deliberately the dispersed oxide as well. Moreover, oxidation of these elements removes them as a precipitation hardener in the alloy. Four techniques, namely the simple mixing, ignition surface coating, internal oxidation and selective reduction, are available to com1

Mechanical Alloying

bine oxide dispersion strengthening and solid solution strengthening in alloy systems containing relatively non-reactive elements. However, all these four techniques were found to be unsuitable for the production of oxide dispersion-strengthened precipitation-hardened nickel-base superalloys [5], mainly due to reactiveness of the chromium, aluminium and titanium present in the alloy. The mixing technique requires interparticle spacing even with a large mechanical reduction during consolidation and subsequent working operations. Powders of this fineness containing γ′ formers such as aluminium and titanium are very reactive, leaving aside the health hazards, because of their high specific surface area and complete or nearly complete oxidation of aluminium and titanium that it can result in. The problem can be overcome to some extent by employing slightly coarser powders and grinding the mixture in a ball mill. To overcome the tendency of fine particles to weld together during the milling, kerosene or fatty acids are usually added. Although lubricants make fine grinding possible, they may severely contaminate the powders and degrade the alloy made from them. The ignition surface coating technique involves mixing the matrix alloy powders with a liquid solution of a salt of reactive metal. This mixture is dried and pulverized, and the powders are heated in an inert or reducing environment converting the salt to a refractory oxide. This technique also produces oxide coated powders which have the same disadvantages as powders made by the simple mechanical mixing technique. In addition, there is a greater contamination problem because of the oxidizing potential of the reaction products of the salt decomposition step. The internal oxidation technique involves oxidation of the alloy powder or thin strip containing a dilute solid solution of the reactive element, in an oxidizing environment at elevated temperatures. It is found experimentally that the particle size of the dispersoid increases with increasing depth of penetration of internal oxidation. Therefore, very fine powders or expensive ultra-thin strips are required to obtain sufficiently fine dispersoid particles. In the case of nickel-based superalloys containing reactive γ′-forming elements, an additional problem would arise. The oxygen potential could not be raised above the extremely low values required to oxidize the γ′−forming elements, which could prohibit the oxidation rate of the desired dispersion-forming elements. The selective reduction process has been used to manufacture commercial dispersion-strengthened materials. It involves production of an intimate mixture of metal oxides and selective reduction of the oxides of the matrix alloy, whilst leaving the dispersoid unreduced. Aluminium and titanium present in the nickel-base superalloy make the reducing 2

Introduction

step impossible with gases, because of the stability of Al2O 3 and TiO 2. These oxides can be reduced by the use of molten alkali and alkaline earth metals. However, this introduces two major new problems, excessive growth of the dispersoid particles, and the necessity to remove the reaction product oxides and carrier agent, usually salt. To cope with these problems associated with the production of oxide dispersion-strengthened nickel-base superalloys J.S. Benjamin (‘Father’ of MA) and his colleagues at the Paul D. Merica Research Laboratory of the International Nickel Company processed a very high purity nickel powder, a fairly coarse chromium powder, a master alloy powder of nickel, aluminium and titanium and a very fine powder of Y 2O3 in a high energy attritor mill in the late 1960s [6]. As a historical note, it should be recorded that at this stage, for the process of mixing/ milling the term ‘mechanical alloying’ was coined by Ewan C. MacQueen, a patent attorney for the International Nickel Company [7]. When such a mixture is mechanically alloyed, the degree to which its various constituents maintain their form depends on their relative hardness and their ability to withstand deformation. Nickel, which is the softest constituent of the mixture, is the cemet that binds the other constituents together. Chromium is somewhat harder and less ductile than nickel, so it tends to form plate-like fragments that are embedded in nickel. The Y 2O 3 disperses along the welds in the composite particles. At the end of the process, the composite particle has a random distribution of oxide dispersoid in a metal-matrix composite. The powder was then consolidated by hot extrusion. The grain size in the extruded bar was very fine. However, to have maximum high temperature strength, the grains should be coarser. For this, the extruded bar was rolled and then annealed at about 1263°C for 30 min, to recrystallize to quite coarse grains. As well as this, they could add elements such as tantalum, molybdenum and tungsten to the nickel-base superalloy, which gave added strength at lower temperatures. 1.2 BENEFITS OF MECHANICAL ALLOYING The MA process has several associated advantages: 1. The homogeneity in fine powder is independent of the initial powder size, which avoids the hazards of fine powders. 2. Fine homogeneous dispersions can be obtained in a particle size of 1 µm or less at a high concentration of alloying elements without occluding air, provided that enough ductile metal powder can be introduced. 3. Grinding times are reduced to 1/10th or even less as compared to that required in a conventional ball or pebble mill. 3

Mechanical Alloying

4. Liquid metal techniques are most convenient and cheaper to develop an alloy. But for the case where it is not possible to get a homogeneous alloy by these techniques, powder metallurgy is adopted. The value of MA becomes apparent when attempts to make an alloy cannot be made by these conventional routes. If the two metals form a solid solution, MA can be used to accomplish the same at lower temperatures. If the two metals are insoluble in solid state, i.e. immiscible solids (e.g. Cu–Fe) or in liquid state i.e. immiscible liquids (e.g. Cu–Pb), an extremely fine dispersion of one of the metals in the other can be achieved. 5. Mechanical alloying represents a cold alloying process, thus it is suitable for hazardous operations. With proper precautions even volatile inflammable materials can be handled safely. Today, MA has been used for developing alloys from immiscible liquids or solids, incongruent melting, intermetallics and metastable phases, and has emerged and developed into a technology capable of providing unique PM materials with consistent properties for high performance applications over a wide range. Various application areas where the MA technology has been utilized are illustrated in Fig.1.1.

Fig.1.1 Application areas of mechanical alloying. 4

Introduction

References 1. 2. 3. 4. 5. 6. 7.

W.D. Coolidge, Proc. Am. Inst. Elec. Eng., 961 (1910). C.S. Smith, Min. and Met., 11, 213 (1930). R. Irman, Tech. Rundschen (Bern), 36, 9 (1949). G.B. Alexander, U.S. Patent 2, 972, 529 (1961). J.S. Benjamin, Met. Trans., 1, 2943 (1970). J.S. Benjamin, Scientific American, 235, 40 (4) (1976). J.S. Benjamin, MPR, 45 (2), 122 (1990).

Questions 1. Which are the various techniques available for preparing dispersion strengthened materials. 2. How is MA helpful in the case of reactive materials? 3. Who invented MA, when and where? 4. Give in brief, a chronological development of MA. 5. Enlist various advantages associated with MA.

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Mechanical Alloying

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MECHANICAL ALLOYING

2.1 ALLOYING MILLS A variety of milling equipment such as attritor mills, vibratory mills, high speed blenders and shakers, planetary mill, and even large diameter conventional ball mills have been used to carry out MA. The milling machine stresses the maximum number of individual particles in a powder mass to undergo plastic deformation or initiate fracture with a minimum of energy. The motion of the milling medium and the charge varies with respect to the movement and trajectories of individual balls, the movement of the mass of balls, and the degree of energy applied to impact, shear attrition and compression forces acting on powder particles. Impact is the instantaneous striking of one object by another. Both objects may be moving or one may be stationary. Attrition is the production of wear debris or particles created by rubbing action between two bodies. This type of milling force is preferred when the materials are friable and exhibit minimal abrasiveness shear consisting of cutting or cleaving of particles, and is usually combined with other types of force. Shear contributes to fracturing by breaking particles into individual pieces, with the creation of a minimal of fines. Compression is the slow application of compressibility forces to a body. A choice between these is likely to be determined by the end results required, as well as the chemical and physical properties of the powder. 2.1.1 Mills in Practice 2.1.1.1 Szegvari attritor mill More than other ball mills, the Szegvari attritor is preferred by most workers because of its operational flexibility. It was invented by Szegvari originally for a comminution/blending machine meant for chemical industries [1]. It consists of a water-cooled stationary vessel with a centrally mounted vertical shaft with impellers radiating from it (Fig.2.1). The shaft is connected to a high-speed, geared motor. The vessel is usually made gas-tight with rubber seals, especially when a controlled atmosphere is required to be maintained. A laboratory attritor is shown in Fig.2.2. When the shaft rotates, arms or ‘lifters’ stir the balls causing them 6

Mechanical Alloying

Fig.2.1 Attritor mill.

Fig.2.2 Laboratory attritor mill (courtesy Union Process). 7

Mechanical Alloying

to lift up and fall back. There is thus a differential movement between the balls and materials being milled, giving a much higher degree of surface contact. The rotating charge of balls and the powder form a vortex at the upper end of the stirring shaft, into which the milling product and balls are drawn. Attritors have many associated advantages as listed below: 1. Milling is achieved by impact and shear forces, and is very intensive because the force restoring the media downward is the weight of all the media above it. 2. The high impact energy allows the use of smaller diameter balls so that powder with a narrow size distribution may be produced. 3. As the greatest milling action is at 2/3 chamber-radius away from the shaft, there is little contamination due to wear of the tank or shaft. The minimal wear of chamber walls ensures long service life. Today, these mills range from the laboratory scale where a few grams of materials are processed to current production units where 1 ton of powder is processed in a 2 m diameter mill which contains more than 1 million balls weighing approximately 10 tons. The alloying process is both mill and powder specific, powders can be in almost any form and can range in diameter from 1–1000 µm [2]. However, these mills have a drawback of relatively low product output (0.5 to 100 kg). They need more power and can be difficult to run and maintain cost effectively. 2.1.1.2 Spex vibratory mill

The use of vibratory ball mills for MA has been promoted predominantly by Kuhn [3,4]. The vibratory ball mill is a long closed tube containing grinding balls and powder. The mill agitates the charge in three mutually perpendicular directions. A schematic of the vibratory mill is shown in Fig.2.3. Impact forces acting on the powder in a vibratory mill are a function of rates of milling, amplitude of vibration and mass of the milling medium. High energy milling forces can be obtained by using high frequency and high amplitude of vibrations. These mills normally operate at frequencies of the order of 200 Hz and a high amplitude of about 12 mm. It is estimated that the maximum accelera-

Fig.2.3 Vibratory mill. 8

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tion of the grinding balls works out to be 24g, where g is the gravitational force [2]. Maximum milling action takes place adjacent to the chamber wall and minimum at the centre, as the motion of the medium decreases from the chamber walls to the centre of the mill tube. The entire charge slowly revolves counterclockwise to the oscillatory vibrations and the high g-forces experienced by the agitated balls leads to intense MA. Vibratory mills utilize small size grinding balls because of higher impact forces, frequencies and acceleration. In fact, it is the most energetic milling device as compared to other mills. These mills are excellent for MA of amounts of up to 4.5 kg or more, depending on the apparent density of the powder. The vibratory mill has, however, not found wide acceptance for bulk production. High speed Spex shaker mills or blenders are ideally suited for laboratory experiment. In fact, some of Benjamin’s experiments were also carried out in a high speed shaker mill. 2.1.1.3 Planetary ball mill Bachin et al [5] carried out MA of dispersion-strengthened, nickel-base superalloys in a centrifugal planetary ball mill. The mechanics of this mill are characterized by the rotational speed of the plate ω p, that of the container relative to the plate ω v, the mass of the charge, the size of the ball, the ball to powder ratio and the radius of the container. A schematic of the planetary ball mill is shown in Fig.2.4. Figure 2.5 shows a laboratory planetary mill. 2.1.1.4 Large diameter ball mills For processing larger batches of powder, however, there has been a trend recently to use conventional horizontal ball mills with larger

Fig.2.4 Schematic of a planetary ball mill. 9

Mechanical Alloying

Fig.2.5 Laboratory planetary mill (courtesy Fritsch GmbH).

diameters (0.5 to 2.5 m) to achieve high energy by rotating it just below the critical speeds ω c (up to 0.9 ω c) [6]. Even though the time required to accomplish MA by these mills is longer compared to attritor mills, the overall economics are favourable. 2.1.1.5 Grinding media

As far as the grinding media are concerned, common practice is to use hardened high carbon–high chromium steel balls (4 to 12 mm diameter), normally specified for ball bearings. Stainless steel balls have also been used. When it is necessary to minimize iron contamination in the charge, balls of tungsten carbide have also been used. When necessary, the balls have been coated with the necessary oxide that was to be dispersed in the composition to be mechanically alloyed [7]. 2.1.2 Improved Mills The commonly used MA mills just discussed are typically inefficient due to the fact that only a small fraction of the expended energy results in the powder deformation, welding and fracture requisites for MA. 10

Mechanical Alloying

Fig.2.6 Regions of the attritor manifesting differing milling actions (Ref. 8, Met. Trans., 24A, 175 (1993)).

These devices are not designed with specific needs of the MA process in mind. This factor has become particularly obvious in recent years, when the dependence of final products on the type of milling device and energy regime used have been repeatedly reported. On these premises a few improved milling devices have been proposed as discussed below. 2.1.2.1 Modified attritor

Attritor efficiency is limited by virtue of its geometrical characteristics and dynamics. The region outside the volume containing the attritor arms is relatively inactive (Fig.2.6) due to a radial ball velocity profile in the mill (Fig.2.7). In fact, a dead zone, containing a disproportionately high volume fraction of powder, exists at the periphery of the mill chamber bottom. However, some regions of the attritor outside the volume containing the arms are characterized by ball velocity gradients which provide a mechanism for mechanical action by sliding and a means for powder circulation within the attritor. The direct impacts, rather than any rolling/sliding events, are primarily responsible for the mechanisms effecting MA [8]. Elimination of the dead zone and a means for facilitating powder circulation within the mill can reduce the process time for an attritor. The following three approaches might lead to improved attritor efficiency [8].

Fig.2.7 Radial ball velocity profile in an operating attritor (Ref.8). 11

Mechanical Alloying

1. The first focuses on the geometrical characteristics of the grinding media; milling with balls having the same diameter results in a closed packed ball array in the attritor volume outside of the core. Such an array results in the predominance of rolling, rather than impact events. Use of balls with two (or more) different diameters would presumably prevent the close packed array from forming, with an attendant increase in mill agitation. The result could be more direct impact events and a lesser relative number of rolling/sliding ones. In fact cinematography has shown that ball mixing disrupts the close packed array assumed by single sized balls in the region of the attritor outside its core. This disruption ostensibly results in a more efficient milling action [9]. 2. The other approaches could facilitate powder removal from the dead zone and/or reduction of this zone’s volume. The bottom of the attritor chamber could be contoured in such a way so as to steer balls and powder up into the more active mill region (Fig.2.8a). 3. An additional arm could be placed on the attritor shaft, as indicated in Fig.2.8b. This (the lowest) arm could be machined into a wedge – as opposed to the current round shape. This would force balls up from the mill bottom into the more active regions. Other balls will fall to fill the volume vacated by balls circulated in this way, thus creating a localized convection of powder and grinding media. 2.1.2.2. Uni-ball mill

Since at the present stage of research very little is known about what milling conditions are to be selected to produce an alloy with a prior

Fig. 2.8 Suggestions for improved attritor design (Ref.8). 12

Mechanical Alloying

Fig.2.9 Schematic of the ball milling with controlled ball movement: 1) rotating cell; 2) balls; 3) magnets (Ref.10, Mat. Sci. and Eng., A134, 1356 (1991)).

chosen structure, there is a need of a simple and versatile device. Ideally, such a mill should offer: – A wide range of milling conditions; – Control of milling parameters. Keeping this in view, the uni-ball mill (universal ball mill) was developed, which is close to these requirements [10]. A schematic of the mill is shown in Fig.2.9. It is a conventional ball mill in which the ball movement during the milling process is confined to the vertical plane by the cell walls and is controlled by an external magnetic field. The intensity and direction of the field can be externally adjusted. By adjusting the spatial and/or temporal profiles of the magnetic field, the ball trajectories, the impact energy and the shearing energy can be varied. Figure 2.10 illustrates the three general patterns of ball movement that can be achieved using this device. In the case illustrated in Fig.2.10a, the magnetic field holds the balls in the bottom part of the cell rotating with the frequency ω c. Friction causes the balls to rotate R in the same direction with the frequency w o = w c c , where Rc is the Rb

Fig.2.10 Modes of operation of the ball mill: a) high energy mode, high rotation frequency; b) low energy mode, low rotation frequency; c) high energy mode two point of equilibrium; d,e) high energy mode, intermediate rotation frequency (Ref.10). 13

Mechanical Alloying

radius of the cell and Rb is the radius of the ball. Periodically, the outer ball on the right-hand side gets released, completes most of the circle being pushed against the cell wall by the centrifugal force and hits the left-most ball of the bottom. With this mode of operation, the powder is worked by both by impact and shearing. The balls may be confined to the bottom part of the cell for the whole time, either by increasing the intensity of the magnetic field or by decreasing the frequency ω c (Fig.2.10b). With this mode, the balls both rotate and oscillate around the equilibrium position of the bottom and the powder is worked mostly by shearing. In the third case, illustrated in Fig.2.10c, the ball movement caused by the centrifugal force can be halted in two opposite directions: at the lowest and highest point inside the cell. The ball trapped by the magnetic attraction in the upper position rotates with frequency ω b and can be released to fall vertically on top of one of the bottom balls. The two colliding balls rotate in opposite directions, which results in a combination of shearing and uniaxial pressure at the surface of contact. Two useful variations of mode (a) are shown in Figs.2.10c and 2.10d. By slowing down the cell rotation frequency ω c a situation may be achieved when the ball released from the bottom is not fully pinned to the wall by the centrifugal force and can either hit one of the bottom balls (Fig.2.10c) or the opposite cell wall (Fig.2.10d). The milling conditions in situations (a), (d) and (c), (e) are similar, respectively. In conventional ball mills, typical impact times are the order of 10 –5 sec and the peak stress can reach 50 kbar. Depending on the collision parameters, the orientation of the impact strain may be quite complex and the temperature rise in the collision region may vary from a few degrees to a few hundred degrees [11,12]. The mechanical energy is utilized most efficiently if the balls collide head-on with maximum possible velocities. Davice et al [12] reported eight head-on collisions out of the total of 2132 collisions per second in a vibratory mill agitated at 1200 rpm. There are only one to two collisions per second in this device, but they are strictly head-on because of the guiding of the ball movement by the magnetic field. The free-fall ball velocity during an impact is 1–5 m/sec (it is increased in the presence of a magnetic field). For comparison, the average relative ball velocity in a vibrating mill varies from 3.9 to 6 m/sec [4], in a planetary mill [13] it remains in the range of 2.5–4.7 m/sec and it is about 0.5 m/sec in attritors [11]. The calculated values of collision time, Hertz radius (refer to Section 9.2.3) and maximum impact stress have been found to be 6.5×10 –5 sec, 4.6×10 –4 and 37 kbar respectively for this mill. All these values are close to the corresponding values quoted for commercial vibrating mills. Thus, the energy released per impact in this mill is not 14

Mechanical Alloying

much different from the characteristics of other devices. The unique feature of the device is the specific ball movement pattern. With every mode of operation this pattern is well defined and highly reproducible. This contrasts with chaotic and unpredicted ball movement characteristics of the other ball-milling devices. Using this improved mill, it could be possible to reproduce both of the amorphization paths: A+B→(AB) am or A+B→(AB) in→(AB)am paths for a Ni–38 at% Zr mixture using two different modes of operation. When the high-energy milling mode is used, the amorphous phase (AB)am is formed directly from the mixture. In contrast to this, when the low energy milling mode (b) is used, the intermetallic phase (AB) in form first, which then slowly transforms into the amorphous phase. Using such a device, it could be possible even to produce metal–metalloid high melting point intermetallics which are believed to require very high energies [14]. To produce fully amorphous Mg 70Zn 30 alloys from the crystalline master alloy having a low crystallization temperature (95°C), a two step ball milling procedure has to be used using the high energy mode (d). First, a steady state mixture of amorphous and crystalline phase is generated, then full amorphization can be achieved by switching to the low energy mode. The high energy mode (d) is not suitable for MA of Al–V, Al–Ti and other aluminium-based and magnesium-based alloys. This mode favours the cold welding process and leads to the formation of large lumps of the milled mixtures. The appropriate milling procedure is to use the low energy mode (b) until the mean grain size decreases to about 0.1 µm and then switch to the intermediate energy mode (e). Such a procedure enables MA of the aluminium based alloys without a PCA (refer to Chapter 4) to control the balance between fracturing and cold welding [15]. 2.2 THE PROCESS Mechanical alloying is normally defined as a dry, high-energy, ball milling process whereby two (or more) elemental powders are blended, cold worked, welded and fragmented repeatedly, resulting in powders with a uniform atom distribution, in stable or metastable phase in a finer microstructure. While applying the technique for dispersoid distribution, it is not really ‘alloying’ due to insolubility of the dispersoids in the matrix. When an alloy powder instead of elemental powders is milled, the process is termed as mechanical grinding (MG) [16]. In this case, the function of MG is to introduce point and lattice defects such as vacancies, interstitials, dislocations, antiphase domain boundaries, etc. to 15

Mechanical Alloying

completely destroy the crystal structure and generate an amorphous material. Thus, the MA process actually promotes particle welding in contrast to the conventional milling practices in which welding is inhibited by the use of liquids and surfactants. Every time two grinding balls collide they trap particles between them. The force of the impact deforms the particles creating atomically clean new surfaces. When the clean surfaces come into contact, they weld together. Since such surfaces readily oxidize, the milling is carried out in an inert atmosphere or vacuum. To facilitate interparticle welding, there must be adequate compressive energy imparted to the grinding medium during milling (hence the high energy mill) and usually the presence of a malleable constituent that could act as a binder for the other constituent and also readily bond with the balls. The other components may include ductile metals, brittle metals, intermetallics, or nonmetals such as carbon, oxides and nitrides. This makes it necessary that milling is done in a dry atmosphere especially for metals of high melting point to promote cold welding. At the early stages in the process, the metal powders are still rather soft and the tendency for them to weld together into larger particles predominates. A broad range of particles develops, with some particles being two to three times larger in diameter than the original ones. As the process continues, the particles get harder and their ability to withstand deformation without fracturing decreases. The smaller particles tend to weld into larger pieces. The large particles, on the other hand, are more likely to incorporate flaws and to break apart when they are struck by the balls. In time, the tendency to weld and the tendency to fracture come into balance, and the size of particles becomes constant within a narrow [17] range (Fig.2.11). The major factors contrib-

Fig.2.11 Balance between welding and fracturing (Ref.17). 16

Mechanical Alloying

uting to this grind limit [18] are: – Increasing resistance to fracture. – Increasing cohesion between particles with decreasing particle size causing agglomeration. – Excessive clearance between impacting surfaces, which is minimized as the ball diameter decreases. – Coating of fine particles on a grinding medium which cushions the particles from impact. – Surface roughness of the grinding medium (highly polished, hard mediums that retain a minimum root mean square surface roughness during milling are most effective). – Bridging of large particles to protect smaller particles in the microbed. Generally, as the alloying proceeds over an extended time, the mean applied stress needed for particle failure increases, while the magnitude of local stresses available to initiate fracture decreases. Although there is little change in size of the particles after the equilibrium is reached, the structure of the particles is steadily refined. The alloying reaches a significant point when the welded layers of a particle can no longer be optically resolved. At this stage, two metals get closely mixed on an atomic level. They have formed a true solid solution rather than a mixture of fine fragments. The welds observed in various systems studied have shown a maximum value of 0.7 µm, but in general this is seen to be far less. At this point, the powder is considered to be adequately processed. It is found that the rate of refinement of the internal structure of

Fig.2.12 Reduction in thickness of the layers within a composite iron–chromium particle during MA (Ref.17). 17

Mechanical Alloying

the particles is roughly logarithmic with processing time (Fig.2.12). As a result, the requirement of fine constituent starting powder is not critical. J.S. Benjamin [19] made the interesting observation that each time a powder particle is trapped between colliding balls, it is plastically deformed sufficiently to reduce its thickness by a factor of two to three times. A simple calculation then leads to the conclusion that for dispersion-strengthening superalloys, starting with a particle size of about 100 µm, only six to eight MA ‘events’ are needed to obtain a desired lamellae thickness within the composite powders of about 0.2 µm. To reduce the inter-lamellae spacing down to five angstroms, or an atomic diameter, requires between eleven and thirteen events. However, the time necessary to cause the required number of collision events to occur is generally in the tens of hours. There is surprisingly little contamination (ppm level) of the powder by the iron in the steel vessel and the steel grinding balls, due to the fact that as the grinding proceeds, the balls and the inner walls of the vessel get coated with a layer of the metals in the mixture [20]. The layer is constantly flaked off the balls, fragmented and rewelded. When a metal is plastically deformed by cold working, most of the mechanical energy of the deformation gets converted into heat (about 5% is stored in the metal raising its internal energy). Heat is also generated by elastic deformation of metal grinding balls and mill chamber walls. The energy expended to overcome the friction between the particles is also translated into heat. Thus, if the temperature of the

Fig.2.13 Typical SEM layered structure of milled powders (Cr–48% Nb elemental mixture) indicating the degree of alloying as a function of milling time: a) 5 hrs; b) 15 hrs; c) 20 hrs; d) 25 hrs (A.H. Clauer and J.J. deBarbadillo (editors). Solid State Powder Processing (1990), p.305 (The Minerals, Metals and Materials Society). 18

Mechanical Alloying

Fig.2.14 Microhardness as a function of processsing time for pure Al + 1.5% NOPCOWAX-22. (Ref.24, Met. Trans., 24A, 513 (1983).

powder rises above a certain point, the cold worked metal particles may undergo recovery and recrystallization. Therefore, water-jacketed milling chambers are usually required for large, high energy vibratory and attritor mills that reach temperatures as high as 200°C [21]. 2.2.1 Process Monitoring The measurement of average flake size does not provide a meaningful criterion for comparing the effect of changes in process parameters, competing alloying processes, and equipment. However, the microstructural changes can reflect the effect of these changes (Fig.2.13) [22,23]. Microhardness measurement is the most meaningful up to the alloying time that produces maximum levels of cold work (Fig.2.14) [24]. X-ray line broadening is sensitive to both the amount of cold work and the refinement in crystallite structure that occurs with continuing kneading and working of the metal well after saturation of cold work [25]. Accordingly, the progress of alloying can be monitored by following microstructural changes, hardness changes and X-ray diffraction. As the absence of sharp X-ray peaks may not always be clear evidence for amorphization, then high resolution electron microscopy [26], scanning electron microscopy [27], Mössbauer spectroscopy [28], and superconducting transition temperature [29] can be used to monitor the process. If peak orientations and/or peak overlapping prevent effective exploration using X-rays, nuclear magnetic resonance spectra [30] can be used. In fact, any measurable property of the material which is affected by MA processing can be chosen to monitor the process. 2.3 FACTORS AFFECTING The progress and the end product of MA is greatly affected by a number of processing parameters, such as mill parameters (impact energy, ball-to-powder ratio (BPR), mill speed, size and size of distri19

Mechanical Alloying

bution of balls, even the shape of impellers in the case of attritor mills), temperature, atmosphere and contamination. 2.3.1 Mill Parameters 2.3.1.1. Impact energy This depends on the specific mill, and the density and size of balls. It is observed that microhardness developed in the MA microstructure is dependent on the impact energy [31,32]. It has also been observed that at high mill energies the degree of crystallization increases and with low energies amorphization occurs [33]. 2.3.1.2 Size of the grinding ball The size of the ball affects size, morphology, recrystallization temperature and enthalpy of the powder produced [34]. As discussed in Section 2.1.2, welding/fracturing events can be enhanced by use of a range of ball sizes, rather than using balls of the same size. 2.3.1.3 Ball–to–powder ratio

Increase in ball–to–powder ratio (BPR) reduces the mean free path of the motion, whilst a low BPR minimizes collision frequency. Thus, impact frequency and total energy consumption per second increase with increasing BPR, while the average impact energy per collision decreases with increasing BPR [31] and minimizes collision frequency. In general, the effective BPR has been found to be in the range 5 to 30. For amorphization cases, it has been found that as the BPR increases the amorphization rate increases sharply, but the contamination, by iron from the milling tools also increases [35]. In general, for amorphization a BPR approaching 100 is frequently used. 2.3.1.4 Speed The milling speed is one of the most important variables to be considered. Very low rotational speeds lead to very long periods of milling (>100 hrs) and a large inhomogeneity in the alloy because of inadequate kinetic energy input, resulting in insufficient localized attribution heat input for alloying. Therefore, an extremely prolonged milling time would presumably be required for homogeneous [20,36] alloying (Fig.2.15). For speeds greater than the optimum, the milling time gets reduced for the same number of revolutions and thus effectiveness of alloying again decreases because of the decrease in the time available for diffusion of the solute. However, very high speeds could lead to excessive heating, high wear of the balls causing contamination from the grinding medium and lower yields. 20

Mechanical Alloying

Fig.2.15 Effectiveness of alloying as a function of attritor RPM for constant number of revolutions Ref.36, Trans. IIM, 39, 596 (1986), The Indian Institute of Metals).

It is conjectured that there may also be an additional factor contributing to the decreased effectiveness of alloying at very high speeds. The horizontal, centrifugal component of the velocity vector is expected to dominate, thus diminishing the vertical vector component of the agitator velocities which is presumably an essential ingredient for ‘homogenizing’ the alloying process, as is the case in a comparable situation of ‘mixing’ by agitators. However, these milling parameters are interdependent. Therefore, it should be adequate to optimize any one of the important parameters for a given set of experimental conditions.

2.3.2 Temperature The ambient temperature of MA and MG is an important parameter which may influence the final structure. The amorphization of the system, for which the heat of mixing is negative (∆H mix< 0) e.g. Ni–Zr, increase in ambient temperature increases amorphization due to the fast diffusion rate of the constituent elements [27,37]. For the system, with positive heat of mixing (∆H mix>0), the effect of ambient temperature starts after a specific period of MA (e.g. Cu–Ta, ambient temperature effect starts after 30 hrs of MA). During this period, the alloying mostly 21

Mechanical Alloying

Fig.2.16 X-ray diffraction spectra for a mixture of Cu30 Ta 70 powders subjected to MA for 100 hrs. (Ref. 37, Mat. Sci. and Eng., A134, 1334 (1991)).

Fig.2.17 Total enthalpy ∆H t vs milling time for the Fe 2 B compound (smooth curve) (Ref.37).

causes the reduction in grain size (≈100 Å) and accumulation of strains. Thus, the interdiffusion begins to act effectively when the average grain size is reduced to a certain minimum. Hence, the retarded temperature effect [37] will be a unique feature observed in a system characterized by a positive ∆H mix (Fig.2.16). The process of MG of intermetallic compounds starts from the lowest free energy. Hence, it involves energetically the up-hill process due to the accumulation of strains and defects, and a reduction in grain size. The total stored energy ∆Ht for the two systems: NiZr2 [38] compound which has been confirmed to amorphize, and Fe 2B compound [37] which does not amorphize, are shown in Fig.2.17. It can be clearly observed that in the case of Fe 2B there is a negligibly small effect of temperature which can be attributed to the lack of formation of the amorphous phase (∆H t levels off at 7.5 kJ/mol, which is smaller than the calculated value of 19 kJ/mol). 2.3.3 Atmosphere Mostly, the MA and MG processes are performed in an inert atmos22

Mechanical Alloying

phere, since a small amount of O 2 or H 2O may have a large influence on the final product. Koch et al [39] have prepared amorphous Ni60Nb40 by MA of the elemental powders in air and helium. The crystallization was found to be different in the two due to different amounts of oxygen content in the alloy. The MA of Ni 60Ti 40 results in different transformation paths when the process is carried out in high-purity argon, air and high-purity nitrogen atmospheres [40]. It is also pointed out that in this case nitrogen has a small influence, while oxygen has a greater effect on MA. 2.3.4 Contamination The type and level of contamination from the wear of the mill chamber and the grinding balls, though small (ppm level), may influence the amorphization transformation path. For example, when MG of Ni–Zr is carried out in an attritor or vibratory mill, the compound transforms directly to the amorphous state. But when the process is carried out in a planetary mill, crystalline intermetallic compound appears as an intermediate product which subsequently transforms to the amorphous state. The difference is attributed to different levels of iron contamination from the different mills [41]. There are cases when the change of grinding medium balls (say from steel to WC) changes the end product. The type of PCA used during MA influences the end product (refer to Chapter 4), which may presumably be due to the level of carbon added in the MA powder by the PCA. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.

W.E. Kuhn and M. Lucky, In: Fine Particles, W.E. Kuhn and J. Ehretsmann (eds). The Electrochem. Soc. Inc. (1974), p.95. G.B. Schaffer and P.G. McCormick, Mat. Forum, 16, 91 (1992). W.E. Kuhn, In: Modern Developments in Powder Metallurgy, V.12, MPIF, Princeton, N.J. (1980), p.195. A.N. Patel and W.E. Kuhn, In: Modern Development in Powder Metallurgy, V.13, MPIF, Princeton, N.J. (1980), p.2750. B.N. Bachin, S.Ya. Kolupaeua, Yu.A. Kustor, A.J. Chernyok and B.V. Schetanov, Poroshk. Metall., 235 (7) (1982), p.44 P.S. Gilman and W.E. Mattson, US Patent 4, 267, 959 (1986). Anonymous, Attrition Mills, MPR, 51 (5), 58 (1996). R.W. Rydin, D. Maurice and T.H. Courtney; Met. Trans., 24A, 175 (1993). T.M. Cook and T.H. Courtney, Met. and Mat. Trans., 26A, 2389 (1995). A. Calka and A.P. Radlinski, Mat. Sci. and Eng., A134, 1356 (1991). D.R. Maurice and T.H. Courtney, Met. Trans., 21A, 289 (1990). R.M. Davice, B. McDermott and C.C. Koch, Met. Trans., 19A, 2867 (1988). J. Eckert, L. Schultz and E. Heustern, J. Appl. Phys., 64, 3224 (1988). A.P. Radlinski and A. Calka, Mat. Sci. and Eng. A134, 1376 (1991). A.P. Radlinski, A. Calka and B.W. Ninham and W.A. Kaczmarek. Mat. Sci. and Eng. A134, 1346 (1991). 23

Mechanical Alloying 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41.

R.B. Schwarz and C.C. Koch, Appl. Phys. Lett., 49, 146 (1986). J.S. Benjamin, Scientific American, 234, 40 (1976). ASM Metals Handbook, Ninth Edition, V.7, 1990, p.56. J.S. Benjamin, Mat. Sci. Forum, 88-90 (1992). E.S. Rao, In: Powder Metallurgy - Recent Advances, V.S. Arunachalam and O.V. Roman (eds), Oxford IBM, New Delhi (1989), p. 27. H. Kimura, M. Kimura and F. Takada, J. Less-Common Metals, 140, 113 (1988). J. S. Benjamin and T.S. Volin, Met. Trans, 5A, 1929 (1974). K.H. Karmer, Powd. Metall. Int., 9, 105 (1977). P.S. Gilman and W.D. Nix, Met. Trans., 12A, 513 (1983). T.K. Wassel and L. Himmel, TACOM-TR-12571 (1981). E. Bonetti, G. Cocco, S. Enzo and G. Valdre, Mat. Sci. and Tech., 6 (12), 1258 (1990). L. Guoxian, W. Erde and W. Zhongren, J. of Mat. Proc. and Tech., 51, 122 (1995). C. Michealsen and E. Hellstern, J. Appl. Phys., 62 (1), 117 (1987). L.M. Di, PI. Loeff and H. Bakker, Mat. Sci. and Eng., A134, 1323 (1991). B.Q. Li and Y.N. Wang, J. Appl. Phys., 75 (3), 1783 (1994). H. Hashimoto and R. Watanabe, Met. Trans. JIP, 31 (3), 219 (1990). E. Gaffet, Mat. Sci. and Eng., 132 (2), 181 (1991). K.B. Gerasimov, J. Mat. Sci., 26 (9), 1296 (1991). P. Yang, J. Jpn. Soc. Powd., Powd Met., 37 (5), 623 (1990). A. Lasonna and M. Magini, Acta Met., 44 (3), 1109 (1996). E.S.B. Rao, R.M. Mallya and D.H. Sastry, Trans. IIM, 39, 596 (1986). C.H. Lee, T. Fukunage and U. Mizutani, Mat. Sci. and Eng.., A134 1334 (1991). C.H. Lee, T. Fukunage and U. Mizutani, Jpn. J. Appl. Phys., 29, 540 (1990). C.C. Koch, D.B. Cavin, C.G. McKamey and J.O. Scarborough, Appl. Phys. Lett., 43 1017 (1983). K.Y. Wang, T.D. Shen, M.X. Quan and J.T. Wang, J. Mat. Sci. Lett., 11, 129 (1992). A.W. Weeber, W.J. Haag, A.J. H. Wester and H. Bakker, J. Less-Common Metals, 140, 119 (1988).

Questions 1 Which are the mills in practice for MA? Which mill is most energetic amongst them? 2. Why, in general, are MA mills energy inefficient. 3. Describe an attritor mill, explaining its salient features. 4. Suggest measures to improve efficiency of an attritor mill. 5. What is a uni-ball mill? Explain specific advantages associated with it? 6. Why is a vibratory mill preferred for amorphization. 7. Give schematic of a planetary mill. 8. Write a brief note on grinding medium used for MA. 9. What advantages are associated with large diameter horizontal ball mills used for mechanical alloying? 10. How is the use of grinding balls of different diameter helpful in improving efficiency of an attritor mill? 11. Explain how the use of a uni-ball mill can avoid use of PCAs in the case of ductile-ductile system? 12. Define MA. How is it different from MG? 24

Variations of Mechanical Alloying

3

VARIATIONS OF MECHANICAL ALLOYING

3.1 REACTION MILLING While most MA is conducted in an inert atmosphere, there can be advantages to MA in a reactive atmosphere. In reaction milling, metal powders react extensively with the milling fluid that is reactive during milling. For example, processing powders in nitrogen atmospheres can result in the formation of nitrides or nitrogen can be trapped in the metal matrix [1,2,3,4], while processing in hydrogen atmospheres may form metastable hydrides [5]. The metal powder is comminuted to a particle size much smaller than the starting one. In general, chemical reactions between the fluid and the powder assist comminution of metal powders by not allowing them to weld together or to the balls/chamber walls. Reaction milling can also be used to obtain the desired dispersoid by introducing an element that reacts during milling or during subsequent heat treatment, or in part during both. For example, MA of Al and C (lamp black or graphite) results in an Al–Al4C 3 composite powder [6], and of Al–B–graphite in an Al–BC composite powder [7]. The process can also be used to produce nanostructured and amorphous materials, if milled for prolonged hours [1], e.g. when iron powder is mechanically alloyed with nitrogen at room temperature, it results in a highly disturbed, non-equilibrium, microfine microstructure with an enhanced nitrogen concentration entrapped in both the metal matrix and the highly disturbed regions. During MA, the impact deforms the particles plastically, creating a new surface. The newly exposed surface is highly reactive to the nitrogen atmosphere and gets adhered to the fresh surface, dissociates and subsequently is incorporated into the matrix when the particles become cold welded. The nitrogen infusion by MA results in a linear increase in nitrogen concentration for up to 100 hrs and a nitrogen level of up to 1.0% in pure iron (the interstitial nitrogen solubility in bcc iron is less than 0.05%) (Fig.3.1). During subsequent consolidation, the matrix nitrogen diffuses rapidly and become trapped in the heavily deformed microstructure and at the subgrain 25

Mechanical Alloying

Fig.3.1 Nitrogen pick-up as a function of MA time (Ref.1).

boundaries. This entrapped nitrogen retards the grain growth and leads to a nanostructural material. 3.2 CRYOMILLING With the recognition of the fact that cryogenic temperatures (which are used to control homologous temperatures) can be used to control the MA process [8], the concept of cryomilling came into existence. In cryomilling, liquid nitrogen is added directly to the milling vessel and the technique has been applied to systems ranging from aluminium to nickel-base superalloys [9,10]. In the case of aluminium, extremely fine Al(ON) particles are reported to be formed in addition to the aluminium oxide present on the surface of the starting powder particles. Cryomilling of Al–50 at% Ti in liquid nitrogen results in the formation and decomposition of B2 TiAl, and the formation of AlN and γ−Al 2O 3 , which impede the grain growth [11]. 3.3 REPEATED ROLLING Mechanical alloying using the repeated rolling technique (Fig.3.2) has been accomplished by several workers [12–14]. The mixture of elemental powders is packed in a stainless tube. The stainless tube which contains the sample powder is pressed to half of its original thickness by a hydraulic press, then cold rolled to a thickness of about 1/20th of the original. The stainless sheath is removed and the rolled sample is packed again in the stainless steel tube of the same size, pressed and rolled. The process is repeated about 30 times. During the repeated rolling process, the thickness of the sample particle d after the n-th cycle of repeated rolling should roughly obey 26

Variations of Mechanical Alloying

Fig.3.2 Schematic of the repeated rolling process (Ref.14: A.H. Clauer and J.J. de Barbadillo (editors). Solid State Powder Processing (1990), p.21. (The Minerals, Metals and Materials Society)).

the equation of kneading, d = d 0 (1/a) n

(3.1)

where d0 is the initial particle diameter and 1/a is the reduction by rolling of one cycle. However, due to changes of the stainless tube after each rolling cycle, the kneading of powders may deviate to some extent from the relationship given in Eq. (3.1). The sharp decrease in reduction of particle size after several initial rolling cycles (Fig.3.3) suggests a change in the kneading mechanism, probably by clipping of particles. P.H. Shingu et al [14,15] have employed this technique for four binary systems; Ag–Fe, Cu–Fe, Ag–Cu and Al–Fe. The heat of mixing for these alloy systems varies from large positive (Ag–Fe) to moderately large negative (Al–Fe). It was shown that due to the simple kneading of elemental powder mixtures by the repeated rolling, metastable alloy phases such as nanocrystalline structure, supersaturated solid solution and amorphous phases can be formed. These results indicate the impact force present in an attritor or ball mill may be effective but not the essential factor for solid state alloying. Thus, this technique can further aid in understanding of the MA process.

Fig.3.3 Plot of the maximum observable thickness of Ag phase vs number of repetitions of rolling process (Ref.19). 27

Mechanical Alloying

Fig.3.4 Flow sheet of the MA process, with two possible routes (Ref.17).

3.4 DOUBLE MECHANICAL ALLOYING In contrast to conventional MA, an alternate process – called double mechanical alloying (DMA) – has been developed [16,17]. Figure 3.4 illustrates the flow sheet of the DMA route. This process production of MA powder involves three main steps. – The first mechanical alloying MA1 step is applied to blend the elemental powders leading to close mixing in a deformed lattice having high dislocation density to facilitate diffusion process which takes place during the second step [18]. – In the second step of the process, the MA1 powder is subjected to high temperature heat treatment under a protective atmosphere promoting intermetallic formation. – In the third step, the heat treated powders are submitted to a second MA process (MA2), where the grain size of the powder being milled is further refined. The technique has been found to be very effective in the case of high temperature aluminium alloys like Al–Fe–Mn and Al–Fe–Ce. High temperature precipitation-hardened Al alloys loose their strength above 200°C due to dissolution of the precipitate and recrystallization of the material. An intermetallic Al6(Fe/Mn) -strengthened Al–Fe–Mn alloy, developed by rapid solidification, was found to have a good combination of strength and ductility [19]. However, direct MA of Al–5% Fe–4% Mn hardly results in the formation of any intermetallic. It results only in refined dispersion of Fe and Mn in the Al matrix [20]. In the DMA technique, a fine distribution of hard Fe and Mn in soft Al results during MA1. During subsequent thermal treatment at 550°C of this powder, a reaction occurs between Al and Fe, and Al and Mn. The first phase formed is metastable Al–Fe. This intermetallic is stabilized by Mn diffusing readily in Al and substituting Fe in the orthorhombic Al 6 (Fe,Mn) intermetallic. The process is limited by the diffusion rate of Mn and Fe in the Al matrix [16]. The size of the intermetallic particles formed ranges from far below 1 µm to a few µm. The properties of the processed powders are shown in Table 3-1. Ex28

Variations of Mechanical Alloying Table 3.1 Properties of Al–5%Fe–4%Mn powders MA1

Thermally treated

MA2

97

97

39

Microhardness (KHN)

1.47

1.31

3.17

Density (g cm–3)

2.83

–

2.85

Carbon ( %)

1.18

–

2.27

Oxygen ( %)

–

–

1

Mean size (µm)

cellent elevated temperature tensile properties are obtained in the consolidated product due to the distribution of high volume fraction of intermetallics and the stabilization of the matrix by inert dispersoids (Al 2 O 3, Al4 C 3). The ductility of the alloy is low, i.e. less than 2% due to the high amount of intermetallics. 3.5 REPEATED POWDER FORGING The highest powder forging has also been used for MA to prepare nonequilibrium and nanostructured materials. Using a novel machine of 500 ton capacity for repeated powder forging, MA has been achieved in Cu–Ag and a range of Al alloys [21]. References 1. 2. 3. 4. 5. 6.

7.

8. 9. 10. 11. 12. 13. 14.

J.C. Rawers and R.C. Doan, Met. and Mater. Trans., 25A 381 (1994). M. Miki, T. Yamasaki, Y. Ogino, Mater. Trans. JIM, 34 (10) (1993) p.952. H. Yasuda, Mater. Sci and Eng., A159, 676 (1994). T.D. Chem, and C.C. Koch, Nanostruct. Mater., 5 (6), 615 (1995). M. Baricco, J. Non-Crystalline Solids, 155, 156 & 527 (1993). H. Danninger, G. Jangg and J. Zbiral, In: Solid State Powder Processing, H. Clauer and J.J. deBarbadillo (eds), The Minerals, Metals and Materials Society, Warrendale, PA (1990), p.241. T. Takahashi and M. Motoyama, Preparation of particle dispersion strengthened aluminium by MA, Presented at Powder Metallurgy World Congress, San Francisco, CA, 21-26 Jan, 1992. J.S. Benjamin and M.J. Bomford, U.S. Patent 3, 816, 080 (1974). R.P. Luton and J. Valone, U.S. Patent 4, 647, 304 (1987). M.J. Luton, Symp. Proc., Mater. Res. Soc. L.E. McCandlish (ed), Pittsburgh, PA, 132 79 (1989). M.J. Luton, Nanostruct. Mater., 5 (6), 631 (1995). M. Atzmon, K.M. Unruh and W.L. Johnson, J. Appl. Phys., 58, 3865 (1985). F. Bordeau, A.R. Yavari and P. Desre, Mat. Sci. and Eng., A197, 129 (1988). P.H. Shingu, K.N. Ishihara, K. Venishi, J. Kuyama, B. Huang and S. Nasu., Ibid ref.6, p.21. 29

Mechanical Alloying 15. 16. 17. 18. 19. 20. 21.

K. Uenishi, K.F. Kobayashi, K.M. Ishihara and P.H. Singhu, Mat. Sci. and Eng., A134, 1342 (1991). P. Le Brun, L. Froyen and L. Delaey, Mat. Sci. and Eng., A157, 79 (1992). X. Niu, P. Le Brun, L. Froyen, C. Peytour and L. Delaey, Powd. Met. Int., 25, (3), 119 (1993). G.B. Schaffer and P.G. McCormick, Met. Trans., 22A 3019 (1991). M.A. Zaidi, J.S. Robinson and T. Sheppard, Mat. Sci. and Tech., 1 737 (1985). P. Le Brun, X. Niu, L. Froyen, B. Munar and L. Delaey, Ibid ref.6, p.273. J. Kihara, Productive Mechanical Alloying by Repeated Powder Forging, Presented at the 1996 World Congress on Powder Metallurgy, Washington DC, WA, June, 1996.

Questions 1. Why are PCAs required during MA? What other benefits are associated with the use of PCAs? 2. What are the commonly used PCAs in MA? 3. How PCAs affect MA powder hardness? 4. Why is the role of PCA most critical during the earliest stage of milling? 5. What factors decide the amount of PCAs being used? 6. Why is degassing necessary for ductile-ductile MA powders? 7. What is the role of oxygen to carbon ratio in organic PCAs. 8. What contamination can PCAs cause? How can this contamination be avoided? 9. In what way can excess use of PCAs be harmful? 10. How use of PCAs affects morphology of powder produced. 11. Give two examples to cite the fact that use of PCAs may lead to structural changes in the powder.

30

Process Control Agents in Mechanical Alloying

4 PROCESS CONTROL AGENTS IN MECHANICAL ALLOYING In the case of ductile–ductile systems (e.g. Al–Mg), it is difficult to use the MA technique because of the excessive cold welding. In such cases, MA results in lumping of the alloyed material, which in turn suppresses the process of alloying. Organic surfactants called ‘process control agents’ (PCAs) are used to achieve the critical balance between cold welding and fracturing, and enhance the process efficiency [1]. The PCAs help in preventing fresh surface contact by giving a surface coating on powders. During the MA process, the PCA is embedded and finely distributed among the layers of flaky powder. The PCAs also help in alleviating the tendency of ductile powder particles towards powder-to-ball/vial welding. A common problem encountered with MA the contamination with elements contained in the mill vial and/or balls. The level of contamination depends on the type of ball mill used, the powder being milled and the milling conditions. The use of a surfactant can also reduce such contamination by at least a factor of 10 [2]. Thus, use of a PCA may be of great practical significance when contamination with surfactant itself does not pose a problem. Benjamin & Bomford used a number of organic compounds including acids, alcohols and ketones [3], for this purpose. Table 4-1 lists the commonly used PCAs during MA.

Table 4.1 Commonly used process control agents Composition ( %) PCA

Chemical Formula C

O

H

N

(COOH) 2 2H 2O

19.1

76.2

4.7

–

CH3OH

37.5

50.0

12.3

–

Stearic Acid

CH3(CH2)16COOH

76.2

11.2

12.6

–

Ethylene bis disteramide (Nopcowax-22 DSP)

C2H 2–2(C18H 36ON)

77.3

5.4

12.4

4.8

Oxalic Acid Methanol

31

Mechanical Alloying

Fig.4.1 Effect of PCA identity and level on welding (elemental Al–2% Cu alloy) (Ref.4, A.H. Clauer and J.J. de Barbadillo (editors). Solid State Powder Processing (1990), p.227).

The following points regarding application of PCAs have been observed experimentally [4]. 1. Approximately the quantity necessary to cover the whole surface area of the final MA mixture with a molecular monolayer of PCA should be added (≈1%) for optimized results. A larger amount may reduce the welding process, hence the powder may become pyrophoric. It is interesting to note that the pyrophoric nature of the powder product with stearic acid added is about 2% (Fig.4.1). As stearic acid has the highest level of carbon of the three PCAs, the phenomenon of pyrophoric is exhibited. 2. The PCA can be in a state of solid, liquid or gas. A fluid PCA reaches the required sites relatively quickly. 3. The efficiency of the PCA is most critical in the early stages of processing (Fig.4.2). 4. The size and shape of the powder produced depends on the ratio of oxygen to carbon present in the PCA. As this increases, the powder size increases and the powder shape becomes more equiaxed (Table 4.2). It appears that hydrogen present in the organic molecular PCAs plays no significant role as a process facilitator. During MA of soft magnetic material (Co–Fe) 75 Si 15B 10 , the use of sodium 1,2 bis (dodecylcarbonyl) ethane-1-sulfonate (cationic) provides a circular shape (≈100 nm) and ammonium dihexadecyl dimethyl-acetate (anionic) provides a needle-like shape (≈100 nm) of MA powders [5]. Table 4.2 Effect of PCA identity on powder size and shape (Ref.4) PCA oxygen/carbon ratio

Powder size

Powder shape

Stearic acid

0.15

Fine

Flaky

Methanol

1.33

Medium

Disc

Oxalic acid

4.00

Coarse

Equiaxed

PCA

32

Process Control Agents in Mechanical Alloying

Fig.4.2 Effect of PCA identity and milling time on welding (elemental Al–2% Cu alloy) (Ref.4). Fig.4.3 (right) Effect of PCA identity and level on PCA powder hardness (elemental Al–2% Cu alloy) (Ref.4).

5. The hardness of the MA powder is also dependent on the type of PCA used (Fig.4.3) in the MA. 6. The use of the PCA may also lead to structural changes in the MA powder, e.g. the use of sodium 1,2 bis (dodecyl carbonyl) ethane1-sulfonate increases the solid solubility of magnesium in aluminium to 13 at %, which is much more than the equilibrium solubility (about 1 at %) in a mixture of Al 70Mg30 (Fig.4.4) [2], (the alloying is judged by the shift of aluminium peaks). The type of PCA used also affects the structure produced, e.g. when Al50Mg 50 is mechanically alloyed for 160 hrs, it produces an amorphous-like structure. After annealing for 5 min at 350°C, this structure transforms into a mixture of two equilibrium phases, β(Al 3Mg2 ) and γ(Al 12 Mg17 ). When different surfactants are used for MA, the final product changes [2] drastically (Fig.4.5). In general, the use of organic PCA contaminates the resulting powder with residual carbon and oxygen [1,6], if a reactive element present

Fig.4.4 X-ray diffractograms of mechanically alloyed elemental mixtures of Al–Mg (helium atmosphere; 80 hrs): a) alloyed without surfactant; b) alloyed with the addition of sodium-1,2 bis (dodecyl carbonyl) ethane-1-sulfonate (Ref.2, Mat.Sci. and Eng., A134, 1346 (1991)). 33

Mechanical Alloying

Fig.4.5 X-ray diffractograms of mechanically alloyed elemental mixtures of Al 50-Mg 50 (helium atmosphere, 160 hrs: a) without surfacant; b) alloyed with addition of sodium-1,2 bis (dodecyl carbonyl) ethane-1 sulfonate; c) dodecyloxycarbonyl sulfosuccinate; d) alloyed with the addition of dodecyldimethylammonium bromide (Ref.2).

in the composition being milled may react with such contaminates [7]. Organic PCAs also form hydrogen during MA which remains in the granulates, necessitating a degassing step. In an effort to reproduce high purity MA material, Gualandi [8] altered this practice by using silicon grease. References 1. 2. 3. 4. 5. 6. 7. 8.

P.S. Gilman and W.D. Nix, Met. Trans., 12A, 813 (1981). A.P. Radlinski, A. Calka, B.W. Ninham and W.A. Kaczmarek, Mat. Sci. and Eng., A134, 1346 (1991). J.S. Benjamin and M.J. Bomford, U.S. Patent 3.816, 080 (1974). J.H. Weber, In: Solid State Powder Processing, A.H. Clauer and J.J. deBarbadillo (eds), The Minerals, Metals and Materials Society, Warrendale, PA (1990), p.227. W.A. Kaczmarek, R. Bramley, A. Calka and B.W. Ninham, In: Conf. Proc. Intermag 90, Brighton, UK, April 1990. F. Faudot, J. Mat. Sci., 28, 2669 (1993). M. Nose, J. Jpn. Soc. Powd. Met., 42 (2), 166 (1995). D. Gualandi and P. Johenson, In: Modern Developments in Powder Metallurgy, H.H. Hausner (ed.) V.3, Plenum Press, New York (1966), p.36.

Questions 1. Describe MA by repeated rolling. 2. What is cryomilling? 3. For which various purposes can reaction milling be used? 4. How is reaction milling able to extend nitrogen solubility in iron? 5. What is DMA? How is it beneficial?

34

Mechanisms in Mechanical Alloying

5

MECHANISMS IN MECHANICAL ALLOYING

The MA technique is being used mainly for three types of processing: alloying, metastable phase formation and activation of chemical reactions. The underlying mechanisms involved in these processes to the extent known today are discussed here. 5.1 ALLOYING To discuss the mechanism of MA, it is convenient to divide the powder charge into three systems: ductile–ductile ductile–brittle brittle–brittle 5.1.1 Ductile–Ductile System When both components are ductile, according to Benjamin and Volin [1], a balance of plastic deformation, cold welding and fracture lead

Fig.5.1 Stages of MA. 35

Mechanical Alloying

to the final microstructure. The alloying occurs in five stages as shown in Fig.5.1. In the first stage, when the particles start getting comminuted, the malleable components are deformed into long lamellae by the impact of the balls while the more friable components are comminuted. This is followed by growth in the number of lamellae due to cold welding; these composite lamellae in the coarser particles have a multilayered, oriented (plate-like) structure. The next stage is associated with the reduction in the aspect ratio of the lamellae and the plate-like coarse particles becoming equiaxed. In the fourth stage, the welding orientation in the composite particles becomes random and convoluted. The final stage is characterized by a narrow particle size distribution and the composition becoming uniform. At this stage, the individual lamellae cannot be resolved in an optical microscope. A saturation level of hardness of the particle is attained and the particles are in the heavily cold worked state. Regarding the atomistic mechanisms that could be responsible for MA, the presence of deformation bands [2], which do not demonstrate in situ recrystallization in MA ductile materials, suggest a collision induced temperature rise (at most a few hundred degrees) is not high enough to account for the estimated diffusion rate of the solute. It is therefore to be surmized that, by some mechanism, the diffusivities are enhanced. One possibility is the role of increased dislocation densities due to severe cold working resulting from MA, which enhance diffusion by providing many sites (tunnels) for diffusion under low activation energy, through pipe diffusion as also through higher vacancy concentrations. As the MA proceeds, plastic deformation thins the lamellae of the powder particles being processed. It decreases the diffusion distance, which further enhances the dissolution of the solute and homogenization. Thus, solute dissolution during MA is facilitated by three factors, namely: – local heating – lattice defects – shortened diffusion distances 5.1.2 Ductile–Brittle System In ductile–brittle systems, as in the case of Ti–Si and Y 2O 3 in superalloys, MA usually results in a fine, homogeneous dispersion of brittle phase in the ductile matrix. The process for this system also follows all five stages of ductile–ductile systems, with the only difference being that the dispersoid present in the constituent is believed to be entrapped along the cold-welded interface [3,4] and its concentration along the weld seams gradually decreases as the welds increase in number and get randomized (Fig.5.2). When the brittle constituent accounts for approximately half the 36

Mechanisms in Mechanical Alloying

Fig.5.2 Dispersoid distribution during MA.

volume fraction, the characteristic layered structure does not develop, instead both the constituents are reduced to nanometer sized crystallites and are evenly distributed throughout the powder [5]. 5.1.3 Brittle–Brittle System In the case of brittle–brittle systems like Ge–Si systems, intermetallic compounds in the Mn–Bi system and amorphous phases from NiZr 2–Ni 11Zr 9 intermetallic, the mechanism of MA is not well understood yet. A granular morphology is observed in the process of MA [6]. There is an initial period of processing which results in very little change in the lattice parameters of the two components. Following this initial period, the lattice parameters shift and eventually converge, indicating complete alloy formation (Fig.5.3). The disappearance of granular-type morphology and the appearance of approximately equiaxed particles, embedded in the major constituent, with a homogeneous microstructure occur at the same milling time where alloying is complete according to the lattice parameter results. A phenomenon of interparticle necking (Fig.5.4) has been observed in MA of such systems, suggesting diffusion under a sufficiently high collision induced temperature rise in contrast to the ductile–ductile system. Such a large temperature rise may be possible in the light of findings of Miller et al, who have showed that a local temperature ranging between 402 to 502°C can be attained in impacted non-metallic crystalline materials such as NaCl and ammonium perchlorate [7] (impact 37

Mechanical Alloying

Fig.5.3 Lattice parameter vs milling time for silicon and germanium powder for the composition Ge–72 at.% Si (Ref.6, Met. Trans., 19A, 2867 (1988). (The Minerals, Metals & Materials Society)). Fig.5.4 (right) SEM micrograph of interparticle 'necking' in Ge–72 at.% Si after milling for 8 hrs (Ref.6, Met. Trans., 19A, 2867 (1988)).

heating due to localized shear deformation, as in metals, may not be possible in brittle systems). Thus, MA of brittle components can be assumed as a thermally activated process. The assumption has also been confirmed by the suppression effect in the alloying process when carried out at cryogenic temperatures using liquid nitrogen [5]. The presence of flat platelets in the mechanically alloyed powder of such materials, in the absence of noticeable slip in these elements at or near room temperature, indicates the temperature-enhanced deformation, provided the temperature approaches 0.5 Tm. However, large scale deformation does not occur as evidenced by the absence of a lamellar microstructure. In the alternative approach to deformation, it is assumed that below 0.5 T m, very little plastic deformation occurs and failure occurs due to defects such as microcracks. This assumption is based on the finding that the brittle materials show a brittle behaviour macroscopically but plastic deformation at microscopic level [8]. Under these conditions along with elevated temperatures, yielding could occur before fracture. Numerous surface asperities are surely present in brittle materials. Upon impact, these asperities of the respective powder particles would penetrate the surface of an adjacent particle. In this way, plastic deformation and bonding might occur. The surface deformation could also be aided by the presence of preferentially active surface dislocation sources. With decreased critical shear stress for dislocation generation at the surface and enhanced mobility provided by somewhat elevated temperatures, plastic flow may be possible, which could lead to the observed MA. 38

Mechanisms in Mechanical Alloying

Plastic flow during MA in the brittle materials can also be understood in the light of the findings of Bridgman [9,10], who showed that the application of superimposed shear or tensile stress in the presence of hydrostatic stress can lead to plastic flow in brittle materials. During the MA process, powder particles are subjected to impact (compressive) forces when entrapped between the balls. Considering that many randomly arranged particles in close contact are involved in a single collision probably triaxial stress states may develop in certain regions of the compact which may have significant hydrostatic components. This could lead to MA via enhanced deformation induced by the hydrostatic stress components. Another possible mechanism involves alloying by material transfer during frictional wear. This is qualitatively consistent with the relatively long processing times observed in brittle–brittle systems compared to ductile–ductile powders. Thus, the mechanisms responsible for material transfer during MA of brittle components may include plastic deformation induced by: – – – –

local temperature rise microdeformation in defect-free volumes surface deformation hydrostatic stresses, or by a frictional wear mechanism.

These temperature induced microdiffusions or deformations may link the mechanisms thought responsible for ductile systems to brittle systems. 5.1.4 Idealness of MA Alloys To identify whether MA alloys are true alloys on an atomic scale or merely a homogeneous mixture at a micron scale, J.S. Benjamin himself and others examined these alloys with the help of magnetic measurements [11,12] and the X-ray technique [13,14], and have demonstrated the true alloying. In contrast to these facts, White and Nix [15] and other investigators found in the Nb–Sn systems that MA was merely an intimate mixture of the constituents and not true alloying. Thus, regarding the ‘true-alloying’ there is no complete agreement. Wassel and Himmel [13] studied this problem in the Cr–Mo system by X-ray diffraction. However, they observed chemical inho-mogeneity in this composition. This could be attributed to their using non-optimized milling parameters. A detailed investigation of the problem was carried out by E.S.B. Rao and his colleagues using nickel–chromium as a model composition. Magnetic coercivity, X-ray line breadth analysis, electron probe 39

Mechanical Alloying

microanalysis (EPMA) and Auger electron spectroscopy (AES) have been used to follow the state of alloying in the attritor milling of Ni20%Cr–2%ThO 2. These studies have revealed that under the conditions of effective milling, magnetic coercivity initially rises to the peak value and drops down, while no such peak is observed if the milling conditions are not conducive for effective alloying (Fig. 5.5). (This experiment is a sensitive test as nickel is a ferromagnetic, but nickel alloy with a chromium content of more than about 8% is paramagnetic). The X-ray peak shifts, after substruction of the strain broadening, also suggest that true alloying could be occurring. The EPMA of mechanically alloyed particles clearly revealed the state of homogeneity of alloying on a micron scale (Fig.5.6). Another interesting set of experiments using AES showed that, with continuous etching of the surface atomic layers of the mechanically alloyed particles, the composition remained homogeneous, confirming the atomic levels of mechanically alloyed particles (Fig.5.7) [16,17]. Thus, MA produces ‘true alloys’ with a homogenous composition provided that the milling conditions are properly optimized. 5.2 METASTABLE PHASE FORMATION 5.2.1 Amorphization Since the discoveries by C.C. Koch et al (1983) of elemental alloy Nb 40Ni 60 by MA [18] and later amorphization of intermetallic by MG (1986) [19], extensive work has gone into the study of amorphous phase formation by MA in the recent past (other methods for solid state amorphization are hydrogen dissolution into a crystalline phase, interdiffusion of two crystalline metals accompanied by a large negative heat of mixing, and irradiation). In the case of MA, structural development takes place due to materials transfer but in the case of MG of elements or intermetallics no materials transfer accompanies the structural changes during ball milling.

Fig.5.5 Magnetic coercivity as a function of milling time at low and high attritor speeds (Ref.17). 40

Mechanisms in Mechanical Alloying

Fig.5.6 The homogeneity as revealed by EPMA line scans of: a) mechanically alloyed elemental nickel and chromium powders; b) atomised Ni–Cr pre-alloyed powder (Ref.17).

Fig.5.7 Homogeneity of composition of mechanically alloyed particles as revealed by the depth profile analysis of AES by sputter etching (Ref.17).

Amorphization by MA Several theories have been suggested for the mechanism of the amorphization phase formation by MA. These include: 1. Liquid quenching (LQ) model This model assumes the occurrence of incipient melting under the impact of colliding balls [20,21] and ultrarapid quenching of the partly molten powder particles. It is natural to think that the severe plastic deformation of the powder particles trapped in ball collisions may locally raise the temperature of the particles sufficiently high to melt, which subsequently solidifies rapidly by heat conduction into the less-deformed and thus cooler regions of the particles. If this mechanism is present, the amorphous phase would form by repeated rapid solidification (RS) processing. However, two observations oppose this simple mechanism. First, the composition range of the alloying does not agree with the range obtained by RS at a cooling rate as high as 10 6 K/sec. Second, although there is no direct measurement of the temperature change in the particles during ball collisions, calculations [19] suggest that the increase is at most a few hundred Kelvin, which is not sufficient for 41

Mechanical Alloying

partial melting of powder particles. 2. Solid state amorphization reaction (SSAR) model This model assumes MA solid state amorphization as in diffusion couples with a high negative heat of mixing and an anomalous high diffusion Table 5.1 Typical list of amorphous alloys formed by MA (Ref.22 and 29).

Alloy A 1–xBx

∆ Hmix (kJ/mole) (kJ/mol)

Atomic size ratio

Hf–Al

–40

1.17

Hf–Cu

–23

1.30

Hf–Ni

–44

1.35

Nb–Cu–Ge

–2

1.14

Nb–Cu–Si

–2

1.14

Nb–Ni

–32

1.18

Pd–Si

–37

–

Sn–Ni

–4

1.31

Sn–Nb

–5

1.11

Ti–Cu

–18

1.15

Ti–Ni

–39

1.19

Zr–Co

–42

1.28

Zr–Fe

–26

1.27

Zr–Ni

–51

1.29

Zr–V

–4

1.19

Cu–W

+33

0.908

Cu–Ta

+3

0.870

Cu–V

+7

0.950

Ni–W

–5

0.855

Cu–Al

–2

0.892

Ta–Al

–46

1.02

Nb–Al

–44

1.03

Ti–Ni–Cu

–

–

Al–Ni–Fe–Gd

–

–

42

Mechanisms in Mechanical Alloying

coefficient [19,21]. The metal–metal amorphous alloys that have been formed by MA are given in Table 5.1. These alloys can be put into two categories: – Metal–metal type (ductile–ductile) alloys consisting of early transitional metal (Ti, Hf, Zr, Nb) and late transitional metal (Fe, Co, Ni, Cu); – Metal–metalloid type (ductile–brittle) alloys like Cu–Nb–(Si, Ge or Sn), Cu–V–(Si or Ge), Nb–(Si or Sn), Pd–Si systems, Pd–N–P and Pt–Ni–P systems. In most of the binary systems, the two elements have the following two important characteristics [22]: – The two elements have a large negative heat of mixing in the amorphous (liquid) state, – The two elements have vastly different atomic sizes. The difference in the atomic sizes is thought to be responsible for the observation that the chemical diffusivity of the two elements in each other and in the amorphous alloys in general differ by several orders of magnitude. Schwarz and Johnson [23] proposed that these characteristics are necessary for the formation of amorphous alloys by interdiffusion at the boundary between two crystalline metals A and B as in the thin film coupled systems, i.e. large negative driving force of the reaction and kinetic suppression of intermetallic formation. The formation of the metastable glassy phase then takes place if the kinetics of transformation are such that the amorphous phase occurs much more quickly than the equilibrium crystalline phase. The shear strain spinoidal transformation makes the crystalline solid unstable and leads to amorphous phase formation at a milling temperature below the glass transition temperature [24]. The plastic deformation, cold welding and fracturing of the particles during MA generate a large concentration of clean metal–metal interfaces, the negative heat of mixing provides a thermodynamic driving force for the reaction enhanced by the excess point, and lattice defects generated by plastic deformation and by the momentary temperature increase in the powder particles trapped in ball collisions. A large difference in chemical diffusivity of the elements in each other and in the amorphous phase favours kinetically the formation of the amorphous alloy in preference to the formation of more stable crystalline intermetallic compounds. The operation of a SSAR during MA is also supported by the formation of the reaction products of MA which are in good agreement with the calculations based on the Phase Diagram method (CALPHAD) [20,25,26]. Figure 5.8a shows a schematic phase diagram for a binary system AB with a negative heat of mixing in the liquid state [21]. 43

Mechanical Alloying

Phases α and β are crystalline primary solid solutions and phase γ is the crystalline intermetallic. Figure 5.8b shows the free energy of α, β and γ, and of the amorphous phase (λ) evaluated at the reaction temperature T r . The free energy of the starting mixture of A and B is located along the straight line joining the free energies of pure crystalline α and β. If the interdiffusion reaction of the A/B interfaces reaches a state of thermodynamic equilibrium, the reaction products would be determined by the common tangents between α, β and γ in Fig.5.8b. However, by choosing a relatively low reaction temperature, the nucleation and growth of the crystalline phase γ is prevented, while still allowing atoms A and B to intermix. This is possible because, owing to their atomic size difference, the chemical diffusivities of A and B in each other and in the amorphous alloy are vastly different. In the absence of γ-phase, the reaction products are determined by the common tangents between phases α, β and λ (Fig.5.8b). These tangents predict five possible reaction products: − a crystalline solid solution, α for 0 < x < x 1; − a two phase product of α (x 1 ) and λ (x 2); − a single-phase amorphous alloy, λ for x 2 < x < x 3; − a two-phase product of λ (x 3) and β (x 4 ); − a crystalline solid solution, β for x 4 < x 0. In the case of AB thin film multilayers, it has been found experimentally that compound nucleation and growth take place when the amorphous layer thickness grows beyond a critical thickness x c. The growth in the critical thickness x c has been correlated with the ratio of diffusivities of the D two species ( xc a B ). A.R. Yavari et al [28] reviewed this for the DA MA case in the light of the fact that the thermodynamic driving force for crystallization disappears with the large concentration gradients is shown to depend on the heat of mixing ∆H mix . Neither ∆H mix < 0 nor D B >>D A are required for amorphization by MA. The critical thickness x c is shown to be ≈2nm for the alloys with ∆H mix=0, such that amorphization by solid-state remains insignificant (in the case of thin films). It is shown, however, that in MA, amorphous layer growth is accompanied by a simultaneous thickness shrinkage due to mechanical deformation. This allows the amorphous interlayer thickness to remain below the critical value x c whilst its volume fraction increases. In another approach to explain the SSAR by MA, Y. Chakk et al [29] developed an atomistic model based on the assumption that amorphization is obtained when a solute atom penetrates into the interstitial sites and distorts the lattice. When such distortions, even when the chemical diffusion is of less significance (i.e. ∆H mix>0 for systems like Cu–Ta), reach some critical value, the long-range order of the lattice is destroyed and the amorphous phase is obtained. It was established that for such a mechanism to be operative, lower and upper limits of the atomic size ratio (ASR) and the minimum solute concentration can be established (refer to Section 9.3). This idea is demonstrated in Fig.5.9. It is divided into four regions. The first region (I) includes systems which are characterized by a low negative enthalpy of mixing and a small difference in atomic radii, as well as pure metallic elements with ∆H mix = 0 and r A/r B = 1. In this region, the systems are not amenable to amorphization by MA. The second region (II) corresponds to a low negative enthalpy of mixing 45

Mechanical Alloying

Fig.5.9 Two-parameter scheme ( ∆H mix vs ASR) to analyze the amorphization tendency of bi-elemental metallic system by MA (Ref.29, Acta Metall. and Mater., 42, 3679 (1994) (Elsevier Science)).

but a large deviation from unity of the ASR. In this region all the systems were successfully amorphized by MA and those systems having low negative or positive heat of mixing could not be amorphized in diffusion couples. The third region (III) contains systems which are characterized by a high negative enthalpy of mixing and the ASR is close to unity. In this region, the systems were amorphized by MA. The forth region (IV) is characterised by a high negative enthalpy of mixing and a large atomic size mismatch. In this region the systems were amorphized by both MA and SSAR of diffusion couples. The ASR criterion should be regarded as a preliminary step for predicting the ability to amorphize a bi-elemental system by MA. The next step was to determine the atomic concentration range where the amorphization can be observed. The maximum strain energy is found to be only about 2% of the total energy stored and reveals that most of the energy of the cold work must be stored in the grain boundaries (of less mobile species) and not in the lattice strain [30]. However, the strain energy has been found to represent about half of the enthalpy of crystallization of the amorphous alloy powder. It therefore can contribute significantly to the amorphization reaction (however, it represents only 8% of that of fusion and 4% of the heat of mixing). Also the rate of amorphization has been found to be maximum when strain is maximum, it is believed that lattice strain plays an important role in the amorphization reaction. Amorphization by MG Amorphous alloy powders have also been prepared by MG of single phase intermetallic powder [30,31], mixtures of intermetallic powders [19] (Ni 32Ti–Ni 45Nb55, NiTi 2–NiNb) and mixtures of intermetallic and elemental powders [32]. (It is surprising that the same amorphous alloy 46

Mechanisms in Mechanical Alloying

powder can be synthesized by ball milling a mixture of elemental powders or by ball milling powders of an intermetallic compound). The MG raises the free energy of the intermetallic to a value equal to or larger than that of the amorphous phase (the reaction from point 3 to point 2, Fig.5.8b). This increase is thought to occur by the chemical disordering of the crystalline lattice, and by the accumulation of point and lattice defects [19,33]. Ball milling causes the lattice of the intermetallic to dilate, which occurs as point and lattice defects are introduced and the lattice is chemically disordered (the chemical disordering replaces strong A–B atomic bonds with weaker A–A and B–B bonds). If the dilation controls the rate of dynamic recovery, then amorphization will be difficult to achieve in crystalline compounds where dynamic recovery limits the defect density and dilation to a value lower than that necessary for the crystal-to-amorphous transformation [34]. (For example, in the MA of nickel and titanium powders to form amorphous Ni 1-x Ti x with (0.3<x 0. T ig is the temperature that is required for the onset of combustion. Both T crit and T ig were found to decrease with milling time [49], which can be attributed to an increase in surface area of the powder system due to a decrease in crystalline size. In addition to the energy dissipated by the collision, the reaction enthalpy also contributes to the temperature rise during such milling, and ultimately causes instability and combustion when the powder temperature reaches the critical value. Figure 5.13 shows a schematic of variation in the critical reaction temperature and local powder temperature maxima with milling time. Curve (a) shows the change in T crit with increasing milling time, whilst curve (c) gives the variation in local powder temperature. The combustion condition is reached after an incu-

Fig.5.13 Schematic showing variation in the critical reaction temperature maxima with milling time (Ref.50, Met. Trans., 23A, 1285 (1992) (The Minerals, Metals and Materials Society)). 52

Mechanisms in Mechanical Alloying

bation milling time tig required for the ignition temperature of the powder mixture Tig to be reduced to a level equal to that achieved locally during collision events. By changing the milling parameters (i.e. milling speed, BPR, ball size), tig the incubation period can be changed but the critical temperature and the ignition temperature remain unaffected for a system [50]. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.

16.

17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29 30 31. 32.

J.S. Benjamin and T.E. Volin, Met. Trans., 5, 1930 (1974). E. Hellstern, H.J. Fecht, C. Garland and W.L. Johnson, MRS Symp. Proc., 132, (1989) 137. J.S. Benjamin and M.J. Bomford, Met. Trans., 7A, 1301 (1977). J.H. Weber, SAMPE Qly., 11 35 (1980). G.B. Schaffer and P.G. McCormick, Met. Trans., 21A, 2787 (1990). R.M. Davis, B. McDermott and C.C. Koch, Met. Trans., 19A, 2867 (1988). P.J. Miller, C.S. Coffey and V.F. Devart, J. Appl. Phys., 54, 913 (1986). H.M. McMillan and N. Gane, J. Appl. Phys., 41, 672 (1970). P.W. Bridgman, Phys. Rev., 48, 825 (1935). P.W. Bridgman, J. Appl. Phys., 18, 246 (1947). J.S. Benjamin, Scientific American, 234, 40 (1976). S.K. Kang and R.C. Benn, Met. Trans., 18A, 747 (1987). T.K. Wassel and I. Himmel, Tr - 12571 (US Army Automotive Command, R and D Control, Michigan) May 1981. B.T. McDermott and C.C. Koch, Scripta Metall., 20, 669 (1986). R.L. White and W. Nix, In: Proc Conf. New Developments and Applications in Composites. K. Wilsdorf and W.C. Harrington Jr. (eds)., The Met. Sco. AIME, N.Y. (1978). E.S.B. Rao, R.M. Mallya. D.H. Sastry and V.S. Arunachalam, In: Proc. Int. Conf, Horizons in Powder Metallurgy, W.A. Kaysser and W.J. Huppman (eds) Verlog Schmid GAAH, Freiburg/Germany (1988), p.1275. E.S.B. Rao, In: Powder Metallurgy - Recent Advances, V.S. Arunachalan and O.V. Roman (eds), Oxford and IBM, New Delhi (1989), p.27. C.C. Koch, O.B. Cavin, C.G. McKamey and Y.O. Scarborough, Appl. Phys. Lett., 43, 1017 (1983). R.B. Schwarz and C.C. Koch, Appl. Phys. Lett., 49, 146 (1986). R.B. Schwarz, R.R. Petrich and C.K. Saw, J. Non-Crystalline Solids, 76, 281 (1985). W. Krauss, C. Politis and P. Weimar, MPR., 43 (4), 231 (1988). R.B. Schwarz and P. Nash, J. of Metals, 91, 27 (1989). R.B. Schwarz and W.L. Johnson , Phys. Rev. Lett., 51, 415 (1983). K.F. Kobayashi, N. Tachibana, P.H. Shingre, J. Mat. Sci., 25, 3149 (1990). N. Saunders and A.P. Miodownik, J. Mater, Res., 1, 38 (1986). Z.H. Yan, M. Oehring and R. Bormann, J. Appl. Phys., 76 (2), 2498 (1992). E. Hellstern, H.J. Fecht, Z. Fu and W.L. Johnson, J. Appl. Phys., 65, 305 (1989). A.R. Yavari and P.J. Derse, Mat. Sci. and Eng., A 134, 1315 (1991). Y. Chakk, S. Berger, B.Z. Weiss and E.B. Levinson, Acta Metall.and Mater., 42 (11), 3679 (1994). A.Ye Yermakov, Ye.Ye. Yurchikov and V.A. Barinov, Phys. Met. Metall., 52, 50 (1981). Ibid ref. 30, 53, 935 (1982). P.Y. Lee, J. Jang and C.C. Koch, In: Solid State Amorphization Transformations, 53

Mechanical Alloying

33. 34. 35. 36. 37. 38. 39.

40. 41. 42. 43 44. 45. 46. 47 48. 49. 50.

R.B. Schwarz and W.L. Johnson (eds), Elsevier, Lausanne (1988), p.28. C.C. Koch and M.S. Kim, J. Phys. (Paris) Colloq., 46, 8-573 (1985). R.B. Schwarz and R.R. Petrich, J. Less-Common Metals, 140, 99 (1988). C.C. Koch, In: Solid State Powder Processing, A.H. Clauer and J.J. deBarbadillo (eds), The Minerals, Metals and Materials Society, Warrendale, PA (1990), p.35. G.J.C. Carpenter and E.M. Schulson, J. Nuclear Mat., 23, 180 (1978). M.S. Kim and C.C. Koch, J. Appl. Phys., 62, 3450 (1987). T.J. Tiainen and R.B. Schwarz, Ibid Ref. 32, p.69. R. Bormann and R, Busch, In: Proc Conf, New Materials by Mechanical Alloying Technique, E. Arzt and L. Schultz (eds), Deutsche Gesellschaft für Metallkunde, Germany (1989), p.73. H.J. Fecht, E. Hellstern, Z.F, and W.L. Johnson, Met. Trans., 21A, 2333 (1990). W.L. Johnson, Prog. Mater. Sci., 30, 81 (1986). W. Gust, B. Predel and K.J. Stenzel, Z. Metallkd, 69, 721 (1978). D.A. Rigney; Annu, Rev. Mater. Sci., 18, 141 (1988). M.L. Trudeau and R, Schultz, Mat. Sci. and Eng., A134, 1361 (1991). M. Oehring and R. Bormann Mat. Sci. and Eng., A134, 1330 (1991). G.B. Schaffer and P.G. McCormick, Met. and Mater. Trans., 21A, 2789 (1990). G.B. Schaffer and P.G. McCormick, Met. and Mater. Trans., 22A, 3019 (1990). G.B. Schaffer and P.G. mcCormick, Scripta Metall, 23, 835 (1989). G.B. Schaffer and P.G. McCormick, J. Mater. Sci. Lett., 9, 1014 (1990). G.B. Schaffer and P.G. McCormick, Met. & Mater. Trans., 23A, 1285 (1992).

Questions 1. Explain the multi-layered structure of MA alloy-powder. 2. Explain the phenomenon of ‘necking’ in brittle-brittle system during MA. 3. How can microdeformation take place in a brittle-brittle system during MA? 4. What are the various factors responsible for MA of a brittle-brittle system? 5. Why was LQ model suitable to explain amorphization by MA? 6. Explain SSAR achieved by MA. How the MA amorphization in systems with positive heat of enthalpy of mixing is caused. 7. Explain role of strain & plastic deformation in MA and MG. 8. Explain amorphization in intermetallic compounds. Why is it not possible to have SSAR in a pure element? 9. Why are nanocrystalline grains considered to be defect free? 10. Explain the role of plastic strain during MA nanocrystallization. 11. Explain MA nanocrystallization in alloys. 12. Differentiate between critical temperature T crit and ignition temperature T ig . 13. How MA helps in promulgating the solid state chemical reactions? 14. What is the effect of milling parameters on solid state chemical reactions? 15. What condition must be satisfied for ignition to start during MA? 54

Energy Transfer and Energy Maps in Mechanical Alloying

6 ENERGY TRANSFER AND ENERGY MAPS IN MECHANICAL ALLOYING A mill is a device which transfers energy to the charge. The mechanism of energy transfer from the milling tools to the milled powder is certainly different with different apparatuses, e.g. in an attritor, an important mechanism of energy transfer is just due to the trapping of the powder between the colliding balls. In a vibratory or planetary mill, the transfer is due to collision of the balls, charged with powder, against the vial walls. For consideration of the energy transfer during the MA process, let us consider the process from fundamentals. What happens when a bare ball falls from a given height (h) on a flat surface of the same material? The ball rebounds to a certain height (h'), giving a rebound yield (n b) = h'/h. Figure 6.1 illustrates a plot of rebound yield vs potential energy of a ball of mass m b that is equal to the kinetic energy of a ball E 0 at the instant of impact ( = m b gh = 1/2 m bv 2 = E 0 ). The yield varies from 0.9 to 0.65 of increasing impact energies [1]. The energy dissipated during an impact E b can be given as E b = (1–n b ) E 0

(6.1)

The dissipated energy gets transformed mainly into heat, resulting

Fig.6.1 Rebound yield of clean balls,η b , vs kinetic energy of the balls, E0 (Ref.1). 55

Mechanical Alloying

in a temperature rise of both the ball and the plate. A small fraction of it is stored in the material as structural disorder. During MA, after a milling period of 10 min to 1 hr, depending on the materials, the balls and chamber walls get coated with a thin powder-layer, and the balls remain ‘glued’ after collision. The collision behaviour approaches that of the elastic type and the energy dissipation becomes minimum (Fig.6.2). In Fig.6.2, the impact energy (x-axis) can be regarded as the intensive milling variable which can be correlated with the impact pressure and temperature rise required to overcome the activation energy barrier. A solid state reaction can occur only if the impact energy is high enough to force the atoms climbing over the activation energy of that reaction. The y-axis represents the quantity of energy released for a collision at any given impact energy. This extensive variable can be correlated to the total quantity of energy, i.e. total time, needed in order to complete a given reaction. Various investigators have developed energy maps for MA/MG for the intermetallic formation (Pd3Si [2], NiAl [3]) and amorphization (NiAl [3], TiNi [4], NiZr [5]). These maps give information about the minimum energy required to start the process and the total energy required for the process to complete. Figure 6.3 illustrates the energy map for the formation of nanocrystalline NiAl intermetallic by MA of Ni and Al elemental powders [3]. It is very clear from Fig.6.3 that a minimum total energy E t = 200 kJ/kg is essential for initiation of the phase formation and 450 kJ/kg for the formation to be complete. It also points out that the completion of the reaction between Ni and Al leading to the formation of single phase NiAl is governed by E t rather than the energy transferred in each impact E b .

Fig.6.2 Energy dissipation during collision, ∆E, vs kinetic energy of the ball E0 (Ref.1). Fig.6.3 (right) Energy map of NiAl formation reaction during MA (Ref.3). 56

Energy Transfer and Energy Maps in Mechanical Alloying

Fig.6.4 Energy map for disordered NiAl formation (Ref.3).

The energy map for the disordering of an ordered NiAl phase is shown in Fig.6.4. It is clear from the energy map that complete disordering of NiAl formed by MA occurs only when the total energy E t crosses 450 kJ/kg. Thus, the accumulation of a critical concentration of defects leads to complete disordering, which is decided by E t not E b. In addition to this, it is interesting to note that the E t required for the completion of NiAl phase formation and for complete disordering are quite close to each other. References 1. 2. 3. 4. 5.

M. Magini, Mat.Sci.Forum, 88-90, 121 (1992). M. Magini and A. Iasomna, Mater. Trans. JIM, 36 123 (1995). J. Joardar et al, In: Recent Advances in Metallurgical Processes, D.H. Sastry, et al (eds), New Age International Publishers, New Delhi (1997), p.647. B.S. Murty et al, Acta Metall., 43, 2443 (1995). E. Gaffet, Mater. Sci. and Eng., A132 181 (1991).

Questions 1. What is the main mode of energy transfer in an attritor mill, a vibratory mill and a planetary mill? 2. Describe phenomena of energy transfer during MA. 3. What are the energy maps used in MA? 4. What is the importance of energy maps in MA? 5. Draw energy maps for MA of (i) Al+Ni (ii) NiAl.

57

Mechanical Alloying

7 CONSOLIDATION OF MECHANICALLY ALLOYED POWDERS 7.1 CONSOLIDATION TECHNIQUES The MA powders have flake-like imperfectly packed layers and rough surfaces which lower their flow rate due to extensive mechanical interlocking. These particles are under heavily worked conditions, which also lowers the shear induced bonding during cold compaction. Thus, conventional methods of consolidation, cold compaction and sintering are generally not suitable for their consolidation. Moreover, due to the layered structure of these powder particles they have a high amount of adsorbed and entrapped gases, which makes degassing, prior to consolidation, an essential step. In the case of nanostructured and amorphous materials, the problems of consolidation are still severe. In the case of nanostructured materials, there is rapid grain growth at elevated temperatures. Grain size thermal stability in these materials can be given by a modified Arrhenius equation [1] d 2 = d 02 + K m t p 2 exp (–Q/RT)

(7.1)

where d0 is the initial grain diameter, K m is the material constant; p, T and t are the pressure, absolute temperature and time used for compaction, respectively, Q is the activation energy and R is the gas constant. In the case of amorphous MA powders, recrystallization at elevated temperatures is a problem [2]. All compaction techniques make use of temperature, pressure and shear. The temperature helps in diffusion bonding and in removing some of the porosity. The pressure provides densification, but does not help bonding. Shear helps to clean the powder surfaces and to bring about interparticle bonding. The various techniques used for the consolidation of MA powders have been compared in Fig.7.1 on the basis of pressure used to achieve compaction and the time–temperature excursion associated with the requirement of complete consolidation (time tem58

Consolidation of Mechanically Alloyed Powders

Fig.7.1 Schematic comparison of techniques for consolidiation of MA powder on the basis of the time-temperature excursion and the stress or pressure used (Ref.3, E. Artz and L. Schultz (editors)), New Materials by Mechanical Alloying (1989), p.143, (Deutsche Geselschaft für Metallkunde)).

Fig.7.2 Schematic comparison of techniques for consolidation of RS and MA powders on the basis of whether shear or pressure forces are applied (Ref.3).

perature parameter = time (sec) × temperature (°C)). Long hold times at high temperatures are useful for obtaining full consolidation by diffusion bonding and by void removal, but leads to a significant loss of metastability [3]. A comparison can also be made on the basis of forces imposed, specifically on the pressure and shear components of these forces (Fig.7.2). The rapid omnidirectional compaction process (ROC) uses a forging press to create high pressures within a thick-walled, hot, solid die. The temperature and forces are sufficient for the die to flow as a fluid (i.e. fluid die), thus creating a nearly hydrostatic pressure [3,4]. Several modifications of the basic process exist to produce shaped parts using simple and cheap fluid dies [4,5]. The choice of technique depends on the material as well as the application as described below. 1. If the material is highly thermally stable, as for example oxide dispersion-strengthened alloys, high temperature techniques can readily be used. 2. For the medium thermally stable materials, for example alloys dispersion strengthened by metallic particles, it is important to restrict the consolidation to sufficiently low temperatures to avoid excessive property degradation. Techniques of high shear, for example extrusion 59

Mechanical Alloying

or forging, are preferred for consolidation. 3. For the highly metastable materials, such as amorphous materials, it is the avoidance of excessive loss of metastability by high temperature excursions which is the most important criterion. Low temperature consolidation techniques such as dynamic compaction, ROC or forging seem to offer the best chances of consolidation. These consolidation techniques provide compacts with nearly full density and have associated advantages but limitations, also which are as follows: 1. Being well characterised and the cheapest process as well as able to provide an isotropic structure, hot extrusion of MA powder is generally preferred [6]. Temperature of the die, extrusion ratio, strain and strain rate decide the grain growth. Higher temperature and lower strain rates tend to allow grain sizes to enlarge gradually, which deteriorates the mechanical properties. 2. Consolidation by forging presents an additional set of necessary controls. 3. In hot isostatic pressing (HIP), the elevated temperatures used during consolidation often result in significant grain growth [1,7]. In addition, the configuration of the final compacts cannot be controlled and is often irregularly shaped. 4. Compaction of MA powders by dynamic compaction techniques like gas gun launchers [8,9] or explosives [10,11] results in minimum grain growth, but due to the high shock pressure of relief waves (a pressure two to three times of the material flow strength is required in these techniques) the compacts often contain fine cracks. Secondly, the need to utilize sufficient energy to bond powder particles which may be suf-

Fig.7.3 Optical micrographs showing dynamically compacted mechanically alloyed Fe40% Ni–14% P–6% B (at.%) amorphous powders: a) at 5 GPa pressure; b) at 4 GPa pressure (Ref.3). 60

Consolidation of Mechanically Alloyed Powders

Fig.7.4 Optical micrographs of MA 760 superalloys: a) as extruded, showing equiaxed grains; b) recrystallized, showing texture and anisotropic grains (Note that both the microstructures show depletion of carbide particles as inhomogeneity) (Courtesy Dr. H.K.D.H. Bhadeshia).

ficient for a temperature rise for localized melting (Fig.7.3a) and recrystallization (Fig.7.3b). Recently, consolidation techniques like cold/warm compaction and sintering [12,13], plasma short activated sintering [14], and liquid phase sintering [15,16] have also been used in the case of thermally stable MA powders. These techniques have also provided compacts with nearly full density and minimum grain growth. However, contamination due to impurities present in the sintering atmosphere is a potential problem. Thus, in consolidation of MA powders, there is difficulty in finding the process ‘window’ where good densification and bonding is obtained, but where the extent of global as well as localised heating and microstructural modification is limited to acceptable values. 61

Mechanical Alloying

7.2 THERMOMECHANICAL TREATMENTS In the case of thermally stable materials, after consolidation several thermomechanical processing steps are necessary to control the grain size and achieve the desired properties in the as-fabricated condition. Two structures are beneficial [17]: 1. Fine equiaxed grains, giving best room temperature strength, fatigue strength and workability. 2. Coarse elongated grains, for high temperature stress-rupture strength, thermal fatigue resistance and corrosion resistance. The fine-grained condition exists in extruded material and, in general, these alloys are not worked at relatively higher values of temperature compensated strain-rate, referred to as the Zener–Holloman parameter [18], which is defined by: z = e& exp (Q/RT)

(7.2)

where e& is the mean strain rate, Q is the activation energy of the ratecontrolling process in the deformation mechanism, R is the universal gas constant and T is absolute temperature. Such strain rates lead to increased substructural strengthening. For example, extruding at a relatively low temperature and high speed produces a high value of z, which can lead to a fine array of subgrains. Hot working at too high value of z can result in sufficient stored energy to cause recrystallization. Coarse elongated grains, e.g. in dispersion-strengthened superalloys, are developed by secondary recrystallization of the materials. Thermomechanical processing produces a high level of stored strain energy [19]. It is a complex matter to control the processing history of the material from the condition of milling, through the temperature, ratio and speed of extrusion, to the details of post extrusion processing, including hot or cold rolling or forging, and the grain coarsening anneal. Achievement of the highest strength attainable by an MA material demands the closest control of production parameters. The recrystallized grains exhibit a texture ({100} or {110} ) and the grains are anisotropic (Fig.7.4). The texture is not only influenced by the recrystallization temperature but also by the temperature gradient conditions [20–22]. Therefore, the desired recrystallized texture must be promoted by controlled nucleation and growth of these very large grains during zone annealing, analogous to what occurs in directional solidification. The development of anisotropic grains in these alloys is attributed to the alignment of oxide particles along the extrusion direction. The alignment is expected to be strongest along the surface regions where the deformation imparted during extrusion is most intense. 62

Consolidation of Mechanically Alloyed Powders

Accordingly, the grains at edges tend to be more anisotropic and less sensitive to heating rates. The core regions are much coarser and sensitive to the heating rate. For randomly dispersed particles, the grain boundary velocity v R is expected to be isotropic. For aligned particles, it should be easier for grain boundary motion to occur parallel to the extrusion direction (velocity v x) and more difficult for motion normal to the extrusion direction v y,z, i.e. v y,z < v R < v x. Since growth along the extrusion direction is less impeded, recrystallization can initiate more readily, thereby explaining the faster kinetics at the surface regions. It is also well established that particles can exert pinning forces. Thus recrystallization in these materials is a complex process. References 1. 2. 3.

4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22.

1. 2. 3. 4.

5.

J. Ravers et al, Met. and Mater. Trans., 27(A), 3126 (1996). C.P. Dogan et al, Nanostruc.Mater., 4(6), 631 (1994). D.G. Morris and M.A. Morris, In: Proc.Conf. New Materials by Mechanical Alloying, E.Artz and L. Schultz (eds), Deutsche Geselschaft für Metallkunde, Germany (1989), p.143. W. Smarsly and W. Bunk, Powd.Met.Int., 17, 63 (1985). R.L. Anderson and J. Groza, MPR, 43 (4), 272 (1988). J.A. Pordoe, Powd. Metall., 22, 22 (1979). C. Suryanarayana, Nanostruc. Mater., 2(5), 527 (1993), D.G. Morris, Mat. Res. Soc. Symp. Proc., 28, 145 (1984). D.G. Morris, Metal Sci., 14 215 (1980). M.A. Meyers et al, J. Metals, 33, 21 (1981). R.N. Wright et al, Adv. Mater. and Proc., 10, 56 (1987). O. Dominguez and J. Bigot, Nanostruc. Mater., 6 (8), 877 (1995). G.E. Fougere et al, Nanostruc. Mater., 5 (2) 127 (1995). K. Kobayashi, J. Jpn. Soc. Powd. Powd. Met., 42 (2), 191 (1995). G. Frommeyer, Paper presented at Euro ’95, Birmingham (1995). Y. Hashimoya, J. Jpn. Soc. Powd. Metall., 42 (2), 185 (1995). M.J. Fleetwood, Mat. Sci. and Techn., 2, 1176 (1986). J.R. Pickens: ASM Metals Handbook, Tenth Edition, 2 200 (1990). R.C. Benn et al, Powd. Metall., 24, 191 (1981). K. Murakani et al, Met. Trans., 24A, 1049 (1993). W. Sha and H.K.D.H. Bhadeshia, Met. and Mater. Trans., 25A, 705 (1994). T.S. Chou and H.K.D.H. Bhadeshia, Met. Trans., 24A, 773 (1993).

Why is ‘conventional compaction & sintering’ not possible in the case of MA powders? What problems are associated in consolidation of MA materials? What is the ROC process? Which dynamic consolidation techniques are being used in case of MA powders. What are basic advantages & limitations of these techniques? Compare, on the basis of pressure & time-temperature excursion parameter, various consolidation techniques used for consolidation of MA materials. 63

Mechanical Alloying

6. 7. 8.

Why thermomechanical treatment is given to metastable MA materials. Why extrusion is normally preferred in the case of MA powder. What disadvantages are associated with HIP when used to consolidate amorphous or nanocrystalline materials?

64

Mechanical Properties of Mechanically Alloyed Materials

8 MECHANICAL PROPERTIES OF MECHANICALLY ALLOYED MATERIALS Mechanical alloying enables the effective superimposition of numerous strengthening factors including: – oxide dispersion – carbide dispersion – fine grain – high dislocation density and substructure – solid solution strengthening Both direct and indirect effects influence the mechanical properties of these materials. The interaction of the second phase particles with dislocations is an example of a direct effect, while grain morphology and texture are both indirect effects. Moreover, the aforementioned five strengthening contributions can be augmented by precipitation strengthening as well as intermetallic dispersion strengthening. Thus, the resulting enhancement in mechanical properties of MA materials is far greater than can be achieved by conventional methods or say by the ‘rule of mixture’. 8.1 TENSILE PROPERTIES The tensile strength of MA materials is found to be relatively proportional to the square of relative density [1] and is consistent with the density–tensile strength relationship for PM materials (Fig.8.1): σ u = k m (ρ measured /ρ full − C) 2

(8.1)

where σ u is the tensile strength, k m is the material constant, ρ is the density and C is a correction factor, equal to green density which represent the state of zero tensile strength. The yield strength of MA materials is found to depend linearly on the inverse of the square root of the grain size (Fig.8.2) (i.e. Hall– Petch relationship), even in the case of nanograin materials [1,2,3]: 65

Mechanical Alloying

Fig.8.1 Tensile stress vs corrected relative density square (Ref.1).

Fig.8.2 Grain size dependence of the yield strength of MA materials (The Int. J. Powd. Met., 24(4), J.S.C. Wang et al, Microstructures and mechanical behaviour of mechanically alloyed nickel aluminide (1988), p.315) (with permission from Prof. C.C. Koch)).

σ y = k m + k 0 d –1/2

(8.2)

where σy is the yield strength, k m is a material constant, k 0 is the grain size coefficient and d is the grain diameter. Thus, the tensile properties of these materials depend strongly on compact density, Eq.8.1, and, to a lesser degree, on grain size. An approximately linear plot even in the case of oxide-dispersed materials, (Fig.8.2), indicates that dispersoids influence the room temperature yield strength through their capability to prevent grain growth during powder consolidation. This tensile behaviour is also consistent with the Orowan mechanism in which the non-shearable oxide particles impede dislocation 66

Mechanical Properties of Mechanically Alloyed Materials

motion by requiring dislocations to bow between particles. The Orowan– Ashby expression [4]

s Or =

0.13 Gb l ln rd / n

b

g

(8.3)

where rd is the dispersoid radius, l is the dispersoid spacing, b is Burger’s vector and G is the shear modulus, gives a value of about 15% for the contribution to the yield strength, compared with the grain size effect, which can be attributed to dispersoid strengthening. The appropriate dispersoid spacing l is given as:

LF p I l = MG J MNH f K

0. 5

-2

OP F 2 I PQ H 3 K

0 .5

rd

(8.4)

where f is the dispersoid volume fraction and rd is the dispersoid particle radius. Equations (8.3) and (8.4) suggest that fine and even distribution of the dispersoid has a stronger effect than its volume. However, in the case of MA intermetallic nanograin materials, the room temperature tensile strength is found to follow the inverse Hall– Petch relationship [5], as in HSLA steels and other nanocrystalline aluminides. This startling behaviour is attributed to the change in the conventional deformation mechanism. Conventionally, in metals of normal grain size the grain boundaries act as barriers to dislocation glide and small grains prevent large dislocation pile-ups. The nanograins are suggested to be essentially dislocation-free [6]. The grain size in these solids is about three orders of magnitude smaller than in superplastic alloys and diffusion rates are appropriately enhanced [7]. It is therefore probable that the mechanism of deformation at room temperature in these materials is similar to the mechanism of superplasticity at elevated temperature like grain boundary sliding and diffusional creep, dislocation glide and creep. However, not all nanocrystalline solids show this inverse Hall–Petch effect [8]. Thus, further experimental work to explain the deformation mechanism explicitly in these materials is required. In general, a simplified relationship exists between ultimate tensile strength and hardness of the metals and alloys [8]:

sy =

Hv km( m -2) (3.0 ± 0.1)

(8.5) 67

Mechanical Alloying

where k m is a constant depending on the material, H v is the diamond pyramid (Vickers) hardness (kg/mm 2 ) and m is the Meyer’s hardness coefficient. The quantity (m–2) is equivalent to the strain hardening coefficient. The assumption of the strain hardening coefficient (m–2) as zero for simplicity in Eq.(8.5), allows an equation of the following form to be used:

sy =

Hv (3.0 ± 0.1)

(8.6)

During MA, powders are subjected to compressive impact forces and undergo severe plastic deformation. During consolidation also these alloys undergo a large plastic deformation (extrusion/rolling/forging, etc). Thus, MA materials undergo extensive plastic deformation before they get the shape, and the strain hardening coefficient can be assumed to be zero [9]. Thus, MA alloys appear to be the ideal candidate for applying Eq.(8.6). In fact, it has been applied for a range of aluminium alloys by I.A. Hawk et al [10] and results have been plotted in Fig.8.3. Equation (8.6) provides a good estimate (with 5% of the actual measured values) of the yield stress for MA alloys not only at room temperature but also after elevated temperature heat treatment due to the stability of fine grains. Equation (8.6) provides a tool for a quick and relatively accurate way to estimate the yield stress of MA alloys for a range of heat treated conditions. It may be specially attractive in the cases of nanostructured materials which are available in limited quantity [10].

Fig.8.3 Plot of yield stress as calculated by Eq.8.6 using hardness values vs the measured tensile yield stress for a series of MA aluminium alloys. (Ref.10, Met. Trans., 19A, 1363 (1988), (The Minerals, Metals and Materials Society)). 68

Mechanical Properties of Mechanically Alloyed Materials

8.2 FRACTURE In general, the ductility and toughness of MA materials is found to be dependent on the morphology of prior particle boundaries [11,12]. The main strengthening factor is found to be the grain size. Elongation is shown to be greater at dynamic strain rates, which is attributed to the initiation of cracks. Dispersoids can harden the matrix by obstructing the motion of dislocations. The dislocation may circumvent the obstacles by passing, cross-clipping, climbing, etc [13,14]. The fracture appears to be dominated by the disaccomodation between the dispersoid particles and the matrix. The fracture path presumably follows crack initiation at the dispersoid oxide/matrix interface and to the extent that dispersoids decorate grain boundaries as well as reside in grain interiors. The fracture paths may be transgranular or intergranular leading to a mixed mode. In the case of nanograin materials, factors such as: – development of secondary cracks; – development of cracks within the nanograin particles and subsequent crack arrest at the particle nanograin–particles and subsequent crack arrest at the particle nanograin–micrograin interface; – crack propagation along the boundary between the nanocrystalline and submicron grain regions increases the fracture toughness [1]. 8.3 CREEP Studies on creep behaviour of various fully dense compacts of MA dispersion-strengthened materials [11,12,15,16] and aluminides [17] have revealed that MA enhances creep resistance and microstructural stability. These improvements are attributed to the fine scale dispersion (~ 30 µm) throughout the matrix, at matrix intermetallic interfaces, and on subgrain boundaries which inhibit coarsening, recovery and recrystallization. The improvement in creep strength, even in the presence of stable fine grain microstructures, in the MA materials may be attributed to the ability of the dispersoids to interact with mobile dislocations over a much wider range of temperature compared to the precipitates present in the conventional alloys [14]. Diffusional creep is generally considered to increase with decreasing grain size. Artz et al [18] proposed that when diffusion distances are small, elevated temperature deformation processes are controlled by the motion of dislocation along grain boundaries rather than the lattice diffusion vacancies. The creep rate is then determined by the resistance offered to the motion of the boundary dislocations by dispersoids situated along grain boundaries in much the same way as Orowan strengthening. The fracture path is presumably dominated by fracture initiation at the dispersoid/matrix interfaces in 69

Mechanical Alloying

these materials. The low creep rates in the case of coarse-grained dispersion-strengthened materials (e.g. thermomechanically treated MA superalloys) can be explained by the modified power-law diffusional creep, which is usually applicable for high temperature alloys, as given below [14]:

e& = km D

LM s N

u

- sth E

OP Q

n

(8.7)

where e& is the creep rate (strain rate), σu is the tensile strength, σ th is the ‘threshold stress’ below which creep rates are considerably smaller, k m is the material constant, D is the diffusivity, E is Young’s modulus and n is the stress exponent. The threshold stress σ th, proposed in modified ‘power-law creep’, Eq.(8.7), is assumed to depend on a parameter k r, describing the relaxation of the dislocation and is proportional to the Orowan stress,

s th = sOr 1 - kr2

(8.8)

8.4 SCC SUSCEPTIBILITY Mechanically alloyed alloys achieve an excellent combination of high strength and superb SCC resistance because of the uniform introduction of microstructural features that either improve SCC resistance, e.g. intermetallic particles, or increase strength without degrading SCC resistance, e.g. oxide and carbide particles, in aluminium alloys. References 1. 2. 3. 4. 5.

6. 7. 8. 9. 10. 11. 12. 13.

J. Rawers et al, Met. and Mater. Trans., 27A, 3126 (1996). T.G. Neil and J. Wadsworth, Scripta Met. and Mater., 25, 955 (1991). K.Y. Wang et al, J. Mat. Sci. Lett., 12 (23), 1818 (1993). W.E. Frazier and M.J. Koczak, In: Dispersion Strengthened Aluminium Alloys, Y.W. Kim and W.M. Griffith (eds), TMS, Warrendale, PA (1988), p.573. A.J. Heron and G.B. Schaffer, In: Advances in Powder Metallurgy and Particulate Materials, V.6, A. Lawley and A. Swanson (compilers), MPIF and APMI, NJ (1993), p.77. S.K. Ganapathi and D.A. Rigney, Scripta Met. and Mater., 24, 1675 (1990). R. Birringer et al, Defect and Diffusion Forum, 59, 17 (1988). J.S.C. Jang and C.C. Koch, Scripta Metall. and Mater., 24, 1599 (1990). J.R. Cahoon et al, Met. Trans., 2, 1979 (1971). I.A. Hawk et al, Met. Trans., 19A, 2363 (1988). C. Zakine, Creep mechanism in mechanically alloyed oxide dispersion strengthened alloys, Presented at 1994 Powder Metallurgy World Conference, Paris, June 1994. H.R. Last and R.K. Garrett, Met. and Mater. Trans., 27A (3), 737 (1996). J.S.C. Wang et al, Inter. J. Powd. Met., 24 (4) 315 (1988). 70

Mechanical Properties of Mechanically Alloyed Materials 14. 15. 16. 17. 18.

R.C. Benn and P.K. Mirchandani, In: New materials by mechanical alloying, E. Arzt and L. Schultz (eds), Deutsche Geselschaft für Metallkunde (1989), p.19. E.L. Erich, Report-AFML-TR-79-4210, US Air Force (1980). A. Lawley and M.J. Koczak, Report No. AFOSR-TR-86-0567 (Oct. 1985). J.W. Pyun and S.I. Kwun, J. Korean Inst. Metals, 33 (6), 814 (1996). E. Arzt et al, Acta Metall, 13, 1977 (1983).

Questions 1. What are the contributing strengthening factors in MA materials? Comment on them. 2. How do tensile properties depend on the grain size in MA materials? 3. How can the inverse Hall-Petch relationship be explained in the case of MA intermetallic nanograin materials? 4. Comment on relationship between yield strength and hardness of MA alloys. 5. Comment on fracture toughness of (i) MA alloys; (ii) Nanograin materials. 6. Explain mechanism of creep in (i) MA fine-grained materials; (ii) MA coarse grained (thermomechanically treated) materials. 7. Why is SCC resistance high in the MA alloys?

71

Mechanical Alloying

9

MODELLING MECHANICAL ALLOYING

9.1 INTRODUCTION Modelling can be simply defined as the representation of a physical entity or a process by another physical or conceptual entity or a process with a view to obtain greater insights into the behaviour of the subject. Models developed for complex processes cannot be expected to be absolutely precise. Rather, they are intended to identify important parameters, define the functional dependence of the process output on process variables and predict results with an acceptable level of precision. One useful result of such process modelling is a considerable reduction in the empirical studies needed to refine a process into a useful engineering tool. Some aspects regarding what is occurring during the MA process have been known qualitatively for some time, but the description of the MA process is complex and manifested as it involves concepts of mechanics, mechanical behaviour, heat flow, thermodynamics and kinetics. In spite of this, modelling of MA has been an avenue of recent interest for further understanding. Though the models for MA are in the early stage of development, various physical/theoretical models developed can be subdivided into mechanistic, atomistic, thermodynamic and kinetic. Mechanistic models deal with powder deformation, coalescence and fracturing, and ball dynamics in a milling device. Atomistic models permit a deeper insight into the physics of non-equilibrium phase formation using the CALPHAD method of thermodynamics data fitting and phase diagram calculations. Kinetic models consider the competitive diffusion-controlled growth kinetics of crystalline phase [1]. In the following section, these are discussed in brief. 9.2 MECHANISTIC MODELS These models deal with the basic process of deformation, coalescence and fragmentation, evaluation of steady state particle size distribution and the milling kinetics. They can be used to predict the milling time required for a particular system, under a given set of experimental con72

Modelling Mechanical Alloying

ditions. Mechanistic models can be classified into two types: local and global: 1. Local modelling describes the various effects (thermal and mechanical) and events (deformation, fracture and welding) that transpire when powder particles are entrapped between two colliding or sliding surfaces. Thus, local modelling is generic in the sense that parameters (relative impact velocity, angle of impact, charge ratio, etc.) affecting the various events are common to all mills, although the values of some of the parameters (e.g. relative impact velocity) are specific to a particular type of mill and its operating conditions. 2. Global modelling is device specific. For example, this type of modelling entails the study of factors such as distribution of impact angles and heterogeneity of powder distribution within the mill, factors which clearly differ from one type of device to another. Several efforts have been made at analysing the ductile–ductile system (to start with) and these are described here. 9.2.1 Deformation, Coalescence and Fracture Regardless of the mill used, the MA is characterised by collisions between tool and powder (ball–powder–ball, ball–powder–vial and ball– powder–impeller arm, etc.) which result in powder fragmentation and coalescence. Such collisions may take place over a range of impact angles, and this geometrical feature might have an important effect on the relative tendencies for coalescence and fragmentation. Figure 9.1 depicts the events possible during MA. Direct seizure (i.e. cold welding taking place without sliding displacement) is favoured by normal impacts, whereas if relative particle displacement precedes such welding, indirect seizure can occur. However, though only a small fraction

Fig.9.1 Various kinds of fragmentation and coalescence events that can be imagined to occur during MA depending upon impact angle (Ref.4, Met. Trans., 21A, 289 (1990)). 73

Mechanical Alloying

of the total number of impacts [3], only the normal collisions are considered. The structure developed in the MA powder depends on the deformation, coalescence and fracture events taking place during the process. Therefore, a model for the deformation, coalescence and fracture of MA events has been developed under the following assumptions [4]: – only the ball–powder–ball collisions are considered as they are greatest in number [2]; – during the processing, a powder coating of about 100 µm thickness gets [2] coated uniformly; – it is assumed that the powder particle shape is a oblong spheroid, though it varies from spherical to flake. This shape is characterised by a ratio (shape factor) of minor and major axis of the spheroid; – it is assumed that these oblong spheroid shaped particles rest on a grinding ball with their major axis parallel to the ball surface. Deformation During the collision, the kinetic energy of the balls is converted into deformation energy during the approach of their centres. In the case of alloying of two or more phase materials, it is taken as the resistance offered by the softer component present. The stress homologous to this energy conversion is applied on the deforming particles. Assuming that the bulk powder particles undergoing collision constitute an ‘individual particle’, the deformation may be expressed as a function within the contact area as

F r IJ a (r ) = R v G HH K

1/ 2

b 0

V

-

r2 Rb

(9.1)

where r is the distance from the centre of contact, R b is the radius of the balls, v 0 is the relative velocity of the balls at impact, ρ is the density of medium balls and H V is the powder hardness. The deformation manifested strain (hardness) can be determined from

e = - ln

LM h - a (r) OP N h Q 0

(9.2)

0

where α(r) is the deformation at a contact radius of r and h 0 is the powder coating thickness. Thus, strain is a radial function of the contact zone. 74

Modelling Mechanical Alloying

Coalescence During plastic deformation, the brittle oxide layer present on the particle surface breaks, the particles flatten in compression, their surface area increases and the underlying metal is progressively exposed. The bonding force F w, acting at the weld that is formed, can be given as F w = A w σu

(9.3)

where σu is the tensile strength of the weld and Aw is the effective new surface area created. If the same metal welds together, σu is the tensile strength of the bulk material. When a cold weld is made between two lamella particles, the weld strength is given by the ‘rule of mixture’. Using the relationship, H v = 3σ u, cf. Eqs. (8.6) and (9.2), one can calculate the bonding force (F w) acting at the weld point which must be more than this for fracture to take place. Fracture Considering only the forging fracture which may develop over several impacts (chances of shear fracture and dynamic fracture are much less during the MA process), the cracks formed will grow radially along the major axes of the particles (Fig.9.2). A crack initiates when the critical tensile strain is attained over a critical length and the crack propagation occurs when the plastic energy release exceeds a value characteristic of the material. Thus, if the particle is sufficiently small so that this length is greater than the particle size, the particle is considered below its comminution limit and will not fracture. The following equation can be used to predict the tensile true fracture strain ε f of a particle ε z (r) = –2ε f

(9.4)

where ε z is the axial strain. When the crack reaches a critical length, determined by the critical value of the J integral, it propagates catastrophically. The value of the J integral is approximately given by [5]

Fig.9.2 Schematic of forging fracture edge cracks formed along the particle circumference and their growth along the particle axis. 75

Mechanical Alloying

J = bsu e a p

F 3s IJ mG H 2s K

m +1

(9.5)

u

where β = 1/ε (σ u /K) m , with K being the strength coefficient and m = 1/n (n is the work-hardening coefficient). The critical J value (J Ic) is related to the critical stress-intensity factor K Ic through J Ic = K 2Ic/ E. The critical crack length a c is obtained for J = J Ic or

Km ac = JIc p m

FG H

2 3 su

IJ K

m+1

(9.6)

Values of J Ic and ac are calculated based on the input values of the material parameters K Ic, E, K, σu and n, where E is Young’s modulus of the powder material. Using the expression for strain as a function of radial position within the powder particle permits determination of the total approach between balls needed in order to exceed some critical strain over a given length. The condition for forging fracture can now be expressed as

F GH

a 2 f 2/ 3 a (r ) = 1 - 1 - c s2 h0 4 Rw

I JK

0 .5

exp( - e c )

(9.7)

where ε c is the critical strain to fracture, (see Eq.9.4), R w is the radius of the weld region and f s is the shape factor. The factor of 4 in the denominator of the last term on the right-hand side of Eq. (9.7) stems from the radial symmetry of the particles: to exceed the critical strain over a c requires that the strain be exceeded over a radial distance one half of a c. The particle shape may affect coalescence and fragmentation events. The shape factor can be expressed in terms of deformation

FG H

fsf = 1 -

a (r ) h0

IJ K

1.5

fsi

(9.8)

where f sf and f si are the final and initial shape factors, respectively.

76

Modelling Mechanical Alloying

9.2.2 Evolution of Particle Size The size and size distribution of the particles change as a result of coalescence and fracture events taking place during the MA. Assuming that respective fracture and welding probabilities are independent of particle size, an expression for the particle size distribution during such ‘fission–fusion’ can be written as [5,6,7]:

z

• dn ( v ) = - a f n ( v ) - a w n( v ) f (v' ) dv' + 0 dt • 1 v a m ( v' )dv' + w n ( v' ) f ( v - v' )dv' + 2a f v v' 2 0

tc

z

z

(9.9)

where α f and αw represent the fracturing and welding probabilities respectively, occurring in a specific impact, t c is the average time between impacts for a particle, n (v) dv is the number of particles having a volume between v and v + dv, and f(v) dv is the corresponding fraction of particles in this volume interval. The first term on the righthand side of Eq.(9.9) represents the removal of a particle by fracture from the volume interval. The second right-hand side term represents particles removal by welding. The third term represents introduction of particles into the volume interval by the fracture of a particle having a volume greater than the one considered. The factor of 2 in this term takes into account that two particles form from one as a result of particle fracturing. The term 1/v' within the integral reflects the assumption that the particle fission distribution is uniform and thus the fission probability is inversely proportional to the number of size intervals the larger particle may break into. The last term represents introduction of a particle into the volume interval by the welding of two particles having the sum of their volume equal to the one considered. The term of 1/2 in this term arises from the convolution integral and is needed to prevent ‘double counting’. There appears to be no analytical solution to such a type of equation. Thus, application of Eq. (9.9) to these processes requires discretisation of the equation. This is done by dividing the population into J equal volume elements, each with an instantaneous population N J. In this formulation, the equivalent of the first term on the right-hand side of Eq.(9.9) is simply –α f N J. The second term is then represented as n- J

-a w

 NJ fI

(9.10)

I =1

77

Mechanical Alloying

Equation (9.10) does not allow particles to weld if the sum of their volumes is larger than a maximum size class. The corresponding discretised form of the third and fourth term on the right-hand side of Eq.(9.9) are n

+2 a f

+

aw 2

NI I -1 I = J +1

Â

(9.11)

J -1

 N J - I fI

(9.12)

I =1

There are two special cases that must be considered during the numerical analysis. The first deals with the smallest size class considered (J = 1). Particles in this size class are not allowed to fracture. In addition, there are no smaller particles that could weld. The equation applicable to this size class is thus written as

LMt N

c

dNJ dt

OP Q

n- J NI N J fI - aw I -1 I = J +1 I =1 n

= 2a f J =1

Â

Â

(9.13)

The second is for the largest size class (J = n). Particles of this size range are not allowed to weld to any other particles. The equation for this size class is

LMt N

c

dN J dt

OP Q

= -a f N J + J =n

aw 2

J -1

 N J - I fI I =1

(9.14)

Fig.9.3 Cu, Nb and Cr particle size distributions obtained from model predictions are compared to those found experimentally for 4 hrs milling (Ref. 8). 78

Modelling Mechanical Alloying

Fig.9.4 Total number of powder particles vs milling time for Cu, Nb, and Cr as observed experimentally (data points) and as predicted by the particle-size evolution model (lines), (Ref.8, Met. Trans., 24A, 2465 (1993)).

The appropriate summation of Eq. (9.10) through (9.14) and stepping through time provide an estimate of the number of particles in a given size class during milling, as would be expected from integrating Eq. (9.9). A bimodel size distribution is predicted by the model and this is found by experiment [8], also in the case of ductile metals like Cu and Nb (Fig.9.3a and 9.3b). However, the model is less satisfactory in the case of brittle metals like Cr (Fig.9.3c). It should be noted that the assumption that fracture and welding probabilities are independent of particle volume and time is not correct. The hardness of powder particles changes concurrent with milling time, with a consequent reduction in welding tendency. In addition, fracture tendencies are also size dependent. Thus, the model is valid for the early stages of milling where fracturing and welding tendencies are approximately constant. The time variation of the total number of particles in the mill can also be predicted by the model and is compared with the experimental results. Such an exercise is illustrated in Fig.9.4, showing only fair agreement between the model and the experimental results, which can be attributed to the respective fracture and welding probabilities used in the model, averaged over the milling time. 9.2.3 Milling Times Since the powder particle aggregates are much smaller than the colliding bodies they are trapped between, it is reasonable to view the surfaces of the colliding bodies as having infinite curvature relative to the powder. The collisions with powder may then be viewed as an upset

Fig.9.5 A cylindrical slug of initial height, h 0 , Hertz radius, r h , entrapped between the colliding surfaces (Ref.4). 79

Mechanical Alloying

forging process between the two parallel plates [9]. We consider a porous metal cylinder, composed of a large number of individual powder particle, of height h 0 and radius r h, as shown in Fig.9.5. The cylinder is impacted between colliding workpieces having an average relative collision velocity v 0. During an impact event, it is assumed that particles confined in the cylinder are subjected to a true deformation strain ε. It is also assumed that the rate of this deformation decreases linearly during the impact. That is, this velocity is v at the initiation of compaction and is zero at the time (τ) at which the impacting workpieces begin to rebound from each other. This model can be used, among other purposes, for estimating the time required for MA. For this first select a ‘unit cell’ or the milling device, i.e. in a SPEX mill, then the unit cell is taken as the mill volume divided by the number of balls within it. In the attritor, the unit cell volume is the average volume ‘belonging’ to each ball. For such a unit cell, the powder volume strain per unit time is επr h2 h 0/t c, where ε is the deformation strain per impact and tc is the time between impact events. To alloy the material, the critical deformation strain ε c, Eq.(9.4), must be generated throughout the total powder volume v p. The processing time t p is the product of V p and ε c divided by the volume strain total per unit time [10] t p = (V p ε c t c)/(επr h2 h 0 )

(9.15)

Processing times thus can be estimated if the ε c, V p, tc, r h2 and h 0 are known. V p is determined experimentally by mill conditions, and ε c can be empirically estimated, assuming the number of collisions to be 3–5 would correspond to a lamellar structure with an initial spacing of 20 µm having its spacing reduced to 1.0–0.1 µm. The estimation of other variables can be explained as follows: Estimation of time between collisions The time between collisions t c can be given by t c = λ/v 0

(9.16)

where λ is the mean free path between collisions. The velocity of balls in an attritor mill can be estimated by considering them as a ‘fluid’. As the agitation speed increases, the balls rise higher against the chamber wall (Fig.9.6). From elementary fluid mechanics, the height of a fluid above the reference line is given by [10,11]: 80

Modelling Mechanical Alloying

Fig.9.6 Agitation of a ball charge in an attritor may be likened to that of a viscous liquid in a spinning tank (Ref.9).

h f = ω 2 R 2 /4g

(9.17)

where ω is the rotational velocity, R is the chamber radius and g is the gravitational constant. Equation (9.17) gives the ball rotational velocity

w=

2 g hf R

d i

0. 5

(9.18)

The ball velocities calculated using Eq.(9.18) are found to be approximately equal to that of the impeller arms. Moreover, balls at the periphery of the attritor mill have the maximum velocity, v = ω R, where R is the chamber radius. The average velocity then can be obtained by weighted average of the velocity of the number of balls. Since the balls move generally in the same direction, the average collision velocity (v 0) is half of this average velocity. On this basis, the collision velocity in an attritor operating under ‘typical’ conditions is 0.53 m/ sec. The mean free path between the collision surfaces is taken as the mean free path between ball surfaces in a loose packed array of balls (packing factor = 0.54). In a Spex mill, it is assumed that the distance between collisions is equal to the vial length. (The assumption is valid when the number of balls in the mill is relatively low, so that ball collisions are infrequent.) The vial accelerates over the first half of each trip and then it begins to decelerate. The balls within the vial continue to travel at the maximum vial velocity. This velocity is twice the average velocity over this range of motion, so the average impact velocity is taken as

81

Mechanical Alloying

v0 =

2x t

(9.19)

where x is the vial length. For a ‘typical’ mill, the average velocity is found to be 3.9 m/sec. These velocities are in good agreement with the studies of Davis et al vis-a-vis a combination of computer simulation and experiment [3]. The impact velocity in the case of a planetary mill can be calculated using the following relationship [12]

LM w F R - R I H 2 K +w =M MM w N 3 v

v0

b

v

p wv

p

OP F RI R P R PP H 2 K Q p

b

v

(9.20)

where ω v and ω p are the angular velocity of the vial and the plate, respectively, and R v and R b are the radius of the vial and the balls, respectively. The impact velocity v 0 in planetary ball mills for ‘typical’ conditions is 11.14 m/sec [13]. This impact velocity is an order of magnitude higher than the impact velocity in an attritor and is a primary reason why processing times in a Spex mill or planetary mill are less than in other MA devices. Estimated collision velocities, times between impacts and mean free paths for common MA devices are given in Table 9.1. For horizontal mills, the diameter of the drum varies from 1 to 2 m and the impact velocity ranges from 4.4 to 6.3 m/sec. These are fairly high. The long processing times required by these mills are related to the longer time between collisions and/or the amount of material impacted per collision.

Table 9.1 Typical impact velocity and impact frequency for common MA devices Impact velocity, vo (m/sec)

Impact frequency 1/tc (sec–1)

Attritor mill

0.53

69

Vibratory mill

3.9

278

Planetary mill

11.24

90.7

Horizontal mill

6.3

–

Device

82

Modelling Mechanical Alloying

Fig.9.7 Collision geometry at maximum compression of a ball impacting on a flat surface.

Estimation of material volume per impact The material volume per impact is a cylinder with the volume r 2h h 0 (Fig.9.5). To know this, one must know the parameters rh and h 0. The radius of the cylinder rh is estimated by considering collisions taking place in the absence of powder. Such collisions can be approximated using Hertz’s theory of impact, which assumes that energy does not dissipate during such impact. The kinetic energy utilized in elastic deformation of the colliding bodies is stored as elastic energy, which is released as the bodies recover. This requirement is met if the relative velocity of the impacting bodies is much less than the speed of sound in the material. Sound velocities (v s ), which can be approximated as (E/ρ) 1/2, where E is the tensile modulus and ρ is the density of the impacting body, are of the order of kilometres per second, i.e. they are about three orders of magnitude greater than the relative velocities of the colliding media. Figure 9.7 shows the collision geometry of maximum compression of a ball impacting on a flat plate having an infinite radius of curvature and presumed mass much greater than that of the ball. If the ball is compressed by a distance δmax, then the impact time (2τ) can be written as

2 t = k0

d max v0

(9.21)

where v 0 is the precollision velocity and k 0 the proportionality constant. The distance, δ max can be expressed in terms of material and machine characteristics (i.e. milling machine, mill parameters) and the relative collision velocity as follows:

b g

2 t = gt v0-0.2 r / E

0. 4

(9.22)

Rb

where Rb is the ball radius, ρ is the ball density, E is the tensile modulus of the colliding media and gt is a parameter which depends on the collision geometry, i.e. ball on ball, ball on plate (value is an order of ten). If the ball and colliding surfaces are of different materials, Eeff will be taken. 83

Mechanical Alloying Table 9-2 Characteristics of Hertzian contacts for various mills (Ref.4) Rb(10–3m)

22τ (10–5sec)

rh(10–4m)

Pmax(109 N/m2)

Stainless steel

2.4

1.62

0.59

2.47

Tungsten steel

2.4

1.38

0.55

5.80

Stainless steel

2.4

1.25

1.99

4.16

Tungsten carbide on stainless

5.6

3.12

4.81

6.18

Stainless steel

6.4

3.08

6.20

3.30

Stainless steel

2.4

1.22

2.10

3.85

Mill

Material

Attritor (ball on ball)

SPEX (ball on flat surface)

SPEX (ball on curved surface)

The radius of contact between the colliding surfaces at maximum compression rh, the Hertz radius, can be expressed as (using the Hertz impact theory):

rh = gr v00.4

F rI H EK

0. 2

Rb

(9.23)

where g r is a geometrical proportionality constant (order of unity). The average pressure developed across the contracting surfaces at maximum compression can be used to determine the elastic strain energy involved in the collision. The pressure is expressed as

Pmax = g p v00.4

F rI H EK

0. 2

E

(9.24)

where gp is also a geometrical constant, values of which range between 0.3 to 0.4. Actual values of impact time, Hertz radius and maximum pressure are provided in Table 9.2 for some common MA devices. It can be noted here that impact times are of the order of 10 –5 sec, Hertz radius of the order of 10 –4 m and impact pressure of 10 9 N/m 2 . When powder is entrapped between the colliding balls, we assume

Fig.9.8 The 'swept' volume model used to determine the volume of powder impacted during MA(The figure is appropriate to a Spex mill) (Ref.9). 84

Modelling Mechanical Alloying

that Hertzian collision is mildly perturbed by the powder deformation, fracture or coalescence processes. Determination of h 0 is somewhat more uncertain. It has been determined using a simple fluid mechanics in conjunction with the geometrical ‘sweeping’ mechanism (Fig.9.8). The cone height in Fig.9.8 is equal to the mean free path between colliding workpieces and has a radius equal to the Hertzian one. It has been estimated that there are of the order of one thousand particles entrapped between workpieces in a single collision. If rh is taken as 2×10 –4 m, a cylinder composed of 1000 spherical particles having a radius of 10 µm and with a 50% packing efficiency would have a height h 0 of 0.7×10 –4 m. Milling efficiency will clearly be greatest when the parameter v 0t/2h 0 is close to unity. The estimated times, Eq. (9.15), are the transient times during which an appropriate steady state particle size distribution is being developed. However, for chemical alloying, a transient time is also needed for diffusional requirements during which alloy particles are developed. These transient periods are not included in these models. Nevertheless, the model provides milling times of the order of the observed ones. Moreover, the predicted relative milling times for the different milling devices are also in order. The success of these mechanistic models depends on effective synthesis of the local and global approaches. For this purpose, two computational programs (MAP1 and MAP2) have been developed [14,15]. Programme MAP1 considers the behaviour of a single species with the option of adding dispersoids. Program MAP2 considers two ductile species (more pertinent), welded to form a third composite lamella (Fig.9.9). However, the prediction of this program is accurate only with a factor of two or so. 9.2.4 Powder Heating The temperature achieved during the collisions is of great significance as the processes like alloying, intermediate phase formation, glass formation, etc., which take place during MA, are profoundly influenced by it. During conventional metal working operations, approximately 95% of the work done is manifested in heat evolution and is so during the MA as well. Assuming the plastic deformation work is entirely converted into an (adiabatic) temperature rise in the impacted powder (per impact), then

z

Vp s u de = Vp Cp DT

(9.25)

85

Mechanical Alloying

where V p is the powder volume impacted and C p is the specific heat of the powder. Equation (9.25) gives the idea about the bulk temperature rise. Depending upon both milling conditions and the powder specific heat this bulk temperature rise can vary from a few degrees to a few hundred degrees [2,3,9]. Lim and Ashby have determined that the difference between the surface temperature (T s) and the bulk temperature (Tb) during such sliding conditions is of the order of 1 K [16]. 9.2.5 Powder Cooling Due to mechanical deformation, the powder temperature increases from a base temperature. If the powders cool to this temperature prior to experiencing an additional impact, the calculated temperature rise can be used directly in analysis of diffusion taking place during MA. As powder particles are small, temperature gradients within them can be considered negligible and the cooling process can be analyzed in terms of Newtonian cooling [9]. The rate of heat removal from the par-

Fig.9.9 MAP2 computional program (Ref.14). 86

Modelling Mechanical Alloying

ticle is the rate of heat transfer to the atmosphere:

-Vp Cp

F dT I = h AbT - T g H dt K t

(9.26)

a

where h t is the heat transfer coefficient, A is the area of the particle exposed to the atmosphere and T a is the ambient temperature. Although particles are disc-like, we assume them to be spherical, and time (t) can be given as

F GH

t = -8.53 ¥ 10 -3 Cp

I ln bT - T g k JK bT - T g

rp2

a

i

(9.27)

a

where k is the particle thermal conductivity, r p the particle radius and T i is the post impact temperature ( = T a + T). The values obtained are of the order of 10–2 sec (for a particle of 50 µm which has been heated to several hundred degrees), which are much less than the time required to have a subsequent collision in the MA mills (Table 9.1). Therefore, it is safe to assume that particles will return to the ambient temperature before being struck again. 9.3 ATOMISTIC MODELS These models deal with the atomic approach to explain amorphization achieved by MA. The structure of a crystalline solid is characterized by a LRO and can be described in terms of the radii of the coordination spheres and lattice symmetry. As distinct from the crystal, an amorphous solid has a short-range order and can be characterized by the peak positions of the pair distribution function. It is postulated that the relative change of the bond length η between the amorphous and crystalline state of a pure metal can be chosen as the criterion for the transition from the crystalline to amorphous state. Thus, η can be expressed in the following way:

FG R IJ - FG R IJ HR K HR K h = FG R IJ HRK a

i

c

i

l

l

i

c

(9.28)

i l

where the superscripts a and c refer to the amorphous and crystalline 87

Mechanical Alloying Table 9.3 Relative changes in bond lengths for fcc, bcc and hcp lattices during amorphization Lattice

η1

η2

η3

fcc bcc hcp

0.027 0.027 0.027

0.185 0.440 0.185

0.110 0.178 0.178

state respectively, i = 1,2,3 are the serial numbers of the co-ordination sphere, R ia is the i-th peak position of the pair distribution function of the amorphous solid, Ric is the radius of the i-th co-ordination sphere of the crystal, and R 1 is the nearest neighbour distance (assuming R l a = R l c). Using Eq. (9.28), the change in bond length due to crystalline to amorphous transition can be calculated and this is given in Table 9.3 [17]. The assumption that the high energy ball milling can result in changes of the bond lengths equal to or greater than the value of η i listed in Table 9.3 would imply that the milling process can lead to crystallineto-amorphous transition of a pure metal. However, an amorphous phase can not be produced by high energy ball milling of only one pure metallic element as at least two elements are needed for successful amorphization by MA [18]. When two elements, A and B, are mechanically alloyed, the impurity atom A can occupy either a substitutional or one of the interstitial positions in the B lattice. The lattice distortion made by an interstitial, which is greater than that of a substitutional atom, can therefore be used as a criterion for the ability to form an amorphous phase in a bi-elemental crystalline system by MA. There should be two limits to this criterion that consist of ASR of the constituent elements. A lower limit below which the impurity atom can easily travel within the host lattice and the distortion of the lattice is negligible, and an upper limit above which the short range order of B atoms is destroyed. Several kinds of interstices exist in lattices, keeping these in view a lower limit of r A/r B, where r B is the radius of the matrix atom and r A is the alloying (impurity) element. An upper limit can be calculated which is used to predict the transformation. The minimum solute concentration CBmin necessary for the amorphization can be given by

CBmin = l 0

vA = l0 vB - vB

1

FG r IJ Hr K B

3

(9.29)

-1

A

88

Modelling Mechanical Alloying

where λ 0 is the atomic size factor, v A and v B are the atomic volumes, respectively. Such models have been successful in predicting the amorphization even in the system with +∆H mix , as discussed in Section 5.2.1. 9.4 THERMODYNAMIC AND KINETIC MODELS The basics of such models have been discussed in Section 5.2.1. The study of these models developed clearly indicate a gap between mechanistic and atomistic models. Thermodynamic and kinetic models are intended to fill this gap. Further progress in this area necessitates the development of novel macrokinetic models linking the plastic strain generation of defects to the mechanism and kinetics (i.e. solid state diffusion and phase transformation) of a metastable phase formation. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10.

11. 12. 13. 14. 15. 16. 17. 18.

B.B. Khina and F.H. Froes, J. Metals, 7, 36 (1996). R.W. Rydin et al, Met. Trans., 24A, 175 (1993). R.M. Devis et al, Met. Trans., 19A, 2867 (1988). D.R. Maurice and T.H. Courtney, Met.Trans., 21A, 289 (1990). B.J.M. Aikin et al, Mat.Sci.and Engn., A147, 229 (1991). J.W. Hilgers et al, Math.Mod., 6, 463 (1985). K.L. Kuttler et al, Appl. Analysis, 19, 75 (1985). B.J.M. Aikin and T.H. Courtney, Met.Trans., 24A, 2465 (1993). D.R. Maurice and T.H. Courtney, Met.Trans., 21A, (2) 289 (1990). T.H. Courtney and D.R. Maurice, In: Solid State Powder Processing, A.H. Clauer and J.J. deBarbadillo (eds), The Minerals, Metals and Materials Society, Warrendale, PA (1990), p.3. V.L. Strater, In: Fluid mechanics, McGraw-Hill, New York (1951), p.25. M. Magini, Mat.Sci.Forum, 88-90, 121 (1992). M. Abdellaoui and E. Gaffet, Acta. Metall. and Mater., 43 (3), 1087 (1995). D. Maurice and T.H. Courtney, Met. and Mater. Trans. A., 26A, 2431 (1995). D. Maurice and T.H. Courtney, Met. and Mater. Trans. A., 26A, 2437 (1995). S.C. Lim and M.F. Ashby, Acta. Metall., 35, 1 (1987). Y. Chakk et al, Acta. Metall., 42 (11), 3676 (1994). P.L. Brun et al, Scripta. Met. and Mater., 26, 1743 (1992).

Questions 1. With the help of a schematic show various possible fracture & coalescence events which may take place during MA. 2. Give a typical account of impact velocity and impact frequency in different MA devices; which mill is considered to be most effective? 3. Describe how the volume of powder impacted in single collision can be estimated. 89

Mechanical Alloying

4. 5. 6. 7. 8. 9. 10. 11.

What is the significance of impact duration? How strain per collision can be estimated. How critical strain to fracture can be estimated. Discuss computational model MAP2 to integrate local & global models. Discuss how impact velocity in the MA devices can be estimated. Describe briefly the thermodynamic model based on CALPHAD and ASR ratio for explaining amorphization by MA. Describe the atomistic models to explain amorphization by MA. Explain amorphization in intermetallic compounds; why is it not possible to have SSAR in a pure element?

90

Joining of Mechanical Alloying Materials

10 JOINING OF MECHANICAL ALLOYING MATERIALS Despite the degassing of MA powders prior to their consolidation, MA products still contain gaseous components which impair the joining of these materials. Various potential joining techniques applicable to MA materials are described here. 10.1 WELDING In general, fusion welding is not suitable for joining MA materials. This technique causes the dispersoid particles to be rejected from the molten metal and the development of a high level of porosity in the weld zones due to excessive degassing (Fig.10.1) [1]. Nevertheless, sound TIG welds suitable for positioning or non load-bearing joints can be made. Fusion welding processes, which minimize the size of the molten zone, produce relatively sound welds. Spot and resistance seam welds with excellent tensile strength can be made. However, it may never be possible to achieve the full stress rupture properties of the base material. Solid state welding techniques, friction welding, diffusion welding, explosive welding and magnetostrictive welding have demonstrated sound weld joints [2]. Figure 10.2 shows the microstructure of a Dispal/Dispal friction weld made by using a continuous drive type friction welding

1 mm

Fig.10.1 Optical micrograph of inert gas tungsten-arc melted Dispal showing porosity in fusion region. (Ref.1, W.A. Kaysser and W.J. Huppmann (editors)), Horizons of Powder Metallurgy, Part II, (1986), p.710 (Verlag-Schmid GmbH)). 91

Mechanical Alloying

500 µm

0.9 µm

30 µm

Fig.10.2 Microstructure of Dispal/Dispal friction weld; (a) and (b) optical micrograph; (c) TEM micrograph of the centre of the friction weld zone (Ref.1).

Fig.10.3 Continuous drive friction welding: basic principle (a); process variables (b).

machine [1]. During the process, one of the components is held stationary, while the other one is rotated at a constant speed (Fig.10.3). As the two surfaces continue to rub against each other, heat is generated at the interface. After a certain weld-time, the rotating component is stopped rapidly and the axial pressure is increased. Thus, the hot material cools under pressure and the weld is consolidated. Due to cooling of the molten pool under pressure, the welds have a fine microstructure without porosity (Fig.10.2), which leads to sound welds. The variation of hardness within the weld zone is shown in Fig.10.4. Hardness in the weld zone remains constant in the case of Dispal/Dispal, but increases in the case of Dispal/steels. The bond strength in the welds is given in Table 10.1. Diffusion welding is performed by giving a suitable heat treatment to the well-aligned polished surfaces. Diffusion welding of polished MA 92

Joining of Mechanical Alloying Materials

Fig.10.4 Hardness of a) Dispal/Dispal and b) Dispal/steel welds (×100) (Ref.1). Table 10.1 Ultimate tensile strength of friction welds At room temperature (MPa)

At 300°C (MPa)

At 400°C (MPa)

Dispal/Dispal*

300

170

150

Dispal/stainless steel

260

175

146

Dispal/carbon steel

210

170

145

Material

* For composition of Dispal refer to Section 12.3

6000 (see Section 12.1) surfaces have been performed by giving a heat treatment at 1180°C for 2 hrs [1]. The welded joint produced had good tensile properties. However, the stress rupture properties in the bond area were inferior to the base material due to fine grains. The problem of fine grains can be overcome by special surface treatment recommended by Helko [3]. However, the welding cycle used by Helko is time consuming and requires well-aligned parallel surfaces as well as a suitable atmosphere to get a sound weld everywhere. 10.2 BRAZING Brazing is applied mainly for attachment of MA components in the aerospace industry. The joint strength is of course limited by the strength of available braze filler materials [2].

93

Mechanical Alloying

Fig.10.5 Optical micrograph of MA 6000, forged bonded and zone-annealed (Ref.5).

10.3 FORGED BONDING A new method called forged bonding has been developed to join the MA materials [4]. In forged bonding, the two surfaces are joined by deforming the bond zone plastically in the fine-grained condition at the same temperature and strain rate as used for bulk forging, followed by subsequent recrystallization of the material to obtain the required grain structure. The grains grow across the bond line, thereby creating a structure with very similar characteristics to the base material. The parameters (temperature, strain rate, strain) have to be controlled carefully. Too low temperatures/too high strain rates lead to cracking and too high temperatures/too low strain rates lead to loss of driving force and hence insufficient recrystallization during the following treatment. Figure 10.5 shows a microstructure of a forged bond in MA 6000 [5]. Grain growth over the bond line and microstructure of the bond region similar to the base can be noticed. The yield stress, ductility, transverse stress rupture strength and HCF properties of the forged bonded specimen are found to be comparable with as-extruded and zone-annealed materials. References 1. 2.

3. 4. 5.

H. Kreve et al, In: Horizons of Powder Metallurgy, W.A. Kaysser and W.J. Huppmann (eds), Verlag Schmid GmbH (1986), p.707. J.J. Fisher and J.H. Weber, In: Proc. of Conf. on Structural Application of Mechanical Alloying, F.H. Froes and J.J. deBarbadillo (eds), ASM International, Materials Park, OH (1990). K.H. Helko, US Patent 3, 787, 748 (1974). R.F. Singer and G.H. Gessinger, Met. Trans., 13A, 1463 (1982). H. Rydstad and R.F. Singer, Ibid Ref.1, p.713.

Questions 1. Why fusion welding techniques are not suitable for MA materials. 2. Give one example where diffusion welding has been used successfully for MA materials. 94

Joining of Mechanical Alloying Materials

3.

What is forged bonding? In which case of MA alloy has it been used successfully?

95

Mechanical Alloying

11 RAPID SOLIDIFICATION AND MECHANICAL ALLOYING Two ‘far from equilibrium’ processing techniques, such as RS and MA have been available for the last three decades to produce materials with properties much superior to those obtained by conventional ingot metallurgy (IM) techniques. The RS technique refers to cooling of metallic melts (most commonly using melt spinning or atomisation techniques) at rates of about 10 4–10 6 K. It is well established that RS can result in one or more of the following constitutional and microstructural modifications [1]: – refinement of grain size and second phase particles or segregations; – extension of solid solubility; – formation of non-equilibrium crystalline or quasi-crystalline intermediate phases; – production of amorphous phases. On the other hand, MA involves repeated welding, fracturing, and rewelding of powder particles in a dry high-energy mill. The process can also lead to the above-mentioned metastable effects, but entirely in the solid state. With either of the techniques, the product obtained (powder only in MA and powder, foil, wire or ribbon in RS) has to be consolidated into larger, more stable forms. To retain all the advantages gained by RS or MA, consolidation should be done at relatively low temperatures. Thus, the challenge is to achieve the necessary consolidation within a temperature window, which is often quite narrow. The consolidation methods described in Chapter 7 are applicable to either of these two types of materials. 11.1 RAPID SOLIDIFICATION VERSUS MECHANICAL ALLOYING A comparison of the RS and MA techniques is given in Table 11.1. 11.2 MECHANICAL ALLOYING OF RAPIDLY SOLIDIFIED POWDERS It is also pertinent to point out that MA of RS powders can provide 96

Rapid Solidification and Mechanical Alloying Table 11.1 Comparison of RS and MA A COMPARISON OF RAPID SOLIDIFICATION AND MECHANICAL ALLOYING 1.

REFINEMENT OF MICROSTRUCTURES High strength alloys can be developed through compositional modification and suitable heat treatments. refined microstructures and higher solute contents, in conjunction with a uniform distribution of fine precipitates during subsequent ageing of the RS alloy, have been found useful for this purpose. For example, addition of about 1% of Fe, Ni, Zr or Mn to conventional 7xxx series aluminium alloys improve their mechanical properties. RS of these alloys results in about 20% increase in YS and UTS without loss of ductility (2). These alloys also possess improved corrosion resistance and stress corrosion cracking resistance. Similarly, addition of eutectoid formers (e.g. Co,Cr, Cu, Fe, Ni etc.) to titanium using RS leads to much higher strengths than normally possible through IM techniques. Table 11.2 lists the room temperature mechanical properties of RS alloys showing the significant improvement achieved through this route. A similar but better improvement in the properties of MA Alloys can be achieved using MA techniques (Table 11–3).

2.

The addition of 1% (about 3.5 at %) Li decreases the density of aluminium by about 3% and increases the elasticity modulus by about 6%. A similar effect is also observed on addition of B to titanium and Li to magnesium. Problem of segregation associated with IM technique of Al–Li can easily be overcome through the RS route. Mechanical properties of the MA AL-Li alloys have an edge over the RS or other PM alloys. however, the Li content should be limited to about 1.3% (13). If added more, it has been observed that fabrication becomes difficult due to high strength of the alloy. Thus, the RS approach allows increased density reduction while MA optimizes mechanical properties.

3.

DISPERSION STRENGTHENING One of the serious difficulties encountered in the precipitation hardened alloys is the coarsening of the precipate on elevated temperature exposure and consequent degradation of the mechanical properties. Thus, development of high-temperature alloys requires very low equilibrium solid solubility of the solute elements and a very low diffusion rate. Addition of Fe, Cr and V transition metals such as to aluminium (4,5), and of rare-earth elements like Er and Nd to titanium (6) by RS techniques are found to be useful in retaining the reasonably high temperature strength levels in these alloys. The transition elements form fine, thermically stable intermetallic dispersoid particles whereas the rare-earth elements form uniformly dispersed oxides in situ. However, sometimes regions near the grain boundaries are devoid in these materials (Fig.11.1). It may be emphasized that MA is a much more powerful technique for dispersion strengthening. Because of the heavy working involved, the matrix grain size is very fine, occasionally reaching the nanometer levels, and further the distribution of dispersoids is much finer and uniform throughout the matrix, at the interface, and on the subgrain boundaries. MA enhances microstructural stability at elevated temperatures (Fig.11.2) , is attributed to the presence of fine scale dispersoids in the microstructure which inhibit coarsening, recovery and recrystallization (7).

4.

EXTENSION OF SOLID SOLUBILITY One of the early advantages realized by RS is the extension of sollid solubility in several alloy systems. This extended solid solubility coupled with the under cooling experience by the melt lead to enhanced solid-solution and fine-grain strengthening. RS can also avoid formation of coarse inter metallics and insolubles in some of the systems. Solid solubility extensions has been achieved in Al–Mg, Al–Si, Al–Cu and Mg–Al alloys by mechanical alloying. Although, these values are significantly higher than those in equilibrium state but lower than values achieved in RS alloys (Table 11.3) (8,9). On the other hand, since MA is carried out completely in the solid state, alloying capabilities are much higher through this route. Addition of magnesium to titanium by conventional methods is difficult becuase magnesium boils before titanium melts; but MA

97

Mechanical Alloying conventional methods is difficult becuase magnesium boils before titanium melts; but MA allow about 3% Mg dissolution in titanium (10). further, MA can produce homogeneous solid solutions in liquid-immiscible (Cu–Pb) and solid–immiscible (Cu–Fe) systems (11). 5.

AMORPHIZATION The free-energy diagrams perdict that the homogeneity range of amorphous alloys formed by RS is usually divided into relatively narrow regimes located near deep eutectics in the liquidus (refer Fig.5.8). In contrast, the composition range of the single phase amorphous alloy prepared by solidstate reaction is wide and continuous, and is located near the centre of the composition range. (Fig.11.3) (12). Mechanical alloying prevents intermetallic phase formation during interdiffusion because of extremely thin layer, thus resulting in a wider glass forming range(13).

6.

MA amorphous phases have a relatively more relaxed disordered structure with a more developed short range ordering and contains a large milling induced stored energy.(14)

7.

Amorphization in many systems like Ni-Pd, and Ni-Ti intermetallic is not possible by RS technique. But MA made the amorphization possible in such cases.(12) The MA amorphouse alloys and rapidly quenched alloys, both exhibit very similar properties like crystallization temperatures and atomic structures.

Table 11.2 Mechanical properties of RS and MA aluminium alloys (Ref.1)

1M 7075 RS 7075(+1/Ni+1/Fe) RS CW67 MA 7075 1M Ti–6A/–4V RS Ti–6A/–4V+1B 1M 2024 RS 2024 MA IN 9021

Temper

0.2% YS (MPa)

UTS (MPa)

%El.

K1c (MPa√m)

T6 Extruded T7 Extruded T6

503 634 580 595

572 717 614 645

10 9 12 3.5

– – 47

– –

848 1090

986 1110

13 2.5

– –

T4 – T4 T6

277 326 500 560

464 542 570 600

22 24 12 12

32 – 30 44

CW67: Al–9.0 Zn–2.5% Mg–1.5% Cu–0.14% Zr–0.1%Ni MA IN 9021: Refer to Section 12.3

synergetic effects and remove many limitations of RS materials as discussed above. Aluminium system Rapidly solidified plus mechanically alloyed (RSMA) aluminium– titanium alloys have been shown to have elevated temperature mechani98

Rapid Solidification and Mechanical Alloying Table 11.3 Extension of solid solubility by MA/RS (Ref.8) Solid solubility (at%) Solvent element

Solute element

Mangesium Aluminum Aluminum Aluminum

Aluminum Magnesium Copper Silicon

Equilibrium ≈ 1 ≈ 1 < 1 ≈ 1

Mechanical alloying

Rapid solidification

3.7 14.1 5.6 4.5

22.4 36.8 17.3 11.1

Fig.11.1 TEM micrographs showing the size and distribution of dispersoids in: a) RS Ti3Al–2% Er alloy;(b) MA Ti3Al–20%Nb–2% Mo–3%V–2% Er (Ref.1, Light Metal Age, 6, 18 (1989)).

Fig.11.2 Grain growth in heat treated ribbon, extruded ribbon-powder, and extruded and heat treated MA powder of Cu–5% Cr (Ref.7, E.Artz and L. Schultz (editors)), New Materials by Mechanical Alloying (1989), p.3 (Deutsche Gesellschaft für Metallkunde).

Fig.11.3 Comparison of the homogeneity ranges of the amorphous phase of Ni–Zr, Co– Zr, and Fe–Zr alloys prepared by MA and RS (Ref.12). 99

Mechanical Alloying

Fig.11.4 Effect of temperature on yield strength of several types of Al–Ti–X alloys (Ref.15).

Fig.11.5 Effect of temperature on tensile elongation of several types of Al–Ti–X alloys (Ref.15).

cal properties superior to those of RS aluminium alloys [15]. The size and distribution of intermetallics are refined through RS and MA introduces a fine distribution of oxide and carbide particles which controls the grain size and enhances grain boundary strength at elevated temperatures. The effect of the RSMA process on the yield strength and elongation for these alloys is shown in Figs.11.4 and 11.5, respectively. The ambient temperature yield strength of melt spun (MS) Al–4%Ti (224 MPa) is 28% greater (Fig.11.4) than that of atomised (AT) Al–4%Ti (178 MPa) due to a faster cooling rate achieved in the case of MS (106 K/sec) than AT (10 4 K/sec), which is believed to result in a higher volume fraction of fine intermetallic particles. Mechanical alloying enhanced both the ambient and elevated temperature strengths of these alloys produced by RS, i.e. gas atomization and melt spinning. AM6 alloy has a room temperature yield strength (325 MPa) 80% greater than that of AT6 alloy (180 MPa). At 300°C, the yield strength of AM6 alloy (196 MPa) is 105% greater than that of AT6 alloy (96 MPa). The 100

Rapid Solidification and Mechanical Alloying

MSMA 185 alloy has an ambient temperature yield strength (526 MPa) 106% greater than that of the RS Al–3%Ti–3% Ce alloy (255 MPa). At 300°C, the yield strength of MSMA 185 alloy (172 MPa) is 25% greater than that of MS Al–3%Ti–3%Ce alloy (137 MPa). Elongation in the MS alloys is found to be 22% (Fig.11.5). The MA, however, reduces ductility. AM6 alloy has an elongation of 9% vis-a-vis alloy AT6 of 22%. Similarly, MA Al–3%Ti–3% Ce has a tensile elongation of 5%, but the MS Al–3%Ti–3% Ce alloy has elongation of 21%. Ostensibly, the reduction in ductility can be attributed to the introduction of oxide and carbide dispersoids during MA. These dispersoids preferentially locate themselves on the grain boundaries and inhibit deformation and limit plastic accommodation. Ternary RS Al–Fe–Ce alloys can be used up to 327°C. However, these alloys lose strength rapidly at high temperatures (see Fig.12.12). To overcome this deficiency, the RS Al–8.4%Fe–3.4%Ce alloy has been mechanically alloyed in a Spex mill to uniformly disperse carbides, oxides and intermetallic phases (metastable Al 10 Fe 2 Ce, and stable Al13Fe3Ce and Al13Fe4 intermetallic crystalline) in the aluminium, which results in the increase in strength and stiffness [16,17]. In another investigation using the RSMA process for Al–4%Cu– 1%Mg–1.5%Fe–0.75%Ce alloy, the strength level could be increased to 435 MPa [18]. Titanium system Two advanced titanium-based alloys, Ti–1% Al–8%V–5% Fe–1% Er and Ti–24%V–10%Cr–5% Er, both with extensive rare earth additions have been developed using the RSMA process. The Ti–1%Al–8%V–5%Fe– 1% Er alloy consolidated directly from a gas atomised powder had coarser beta grains of about 30 µm. A MA of up to 40 hours [19] did not yield a totally uniform structure, but the majority of the structure in the consolidated alloy consisted of submicron grains with dispersion of 30 to 50 nm in size. The fine-grain structure was found to be stable with little grain coarsening, and a little drop in hardness was observed on solution treatment at 675°C. The Ti–24%V–10%Cr–5%Er alloy in the as-atomised condition was badly segregated with large chunks of free Er. The microstructure of the consolidated RSMA alloy consisted of both very fine grained regions and some coarser grains, with dispersion in fine grained regions of less than 10 nm in size. The RSMA alloy had higher hardness (437 in comparison with 304 KHN) due to the grain refinement and dispersoids. Thus, fine dispersions and stable grain refinements are possible by 101

Mechanical Alloying

adopting a RSMA process which leads to improvement in the properties of RS materials.

References 1. 2. 3. 4. 5.

6. 7.

8. 9. 10. 11. 12. 13. 14. 15.

16.

17. 18. 19.

C. Suryanarayana and F.H. Froes, Light Metal Age, 6, 18 (1989). P.K. Domalavage et al, Met. Trans., 14, 1599 (1983). W.E. Quist et al, In: Aluminium-Lithium Alloys, C. Baker et al (eds), The Institution of London (1986), p.625. P.S. Gilman et al, Industrial Heating, 56 (2), 30 (1989). Y.M. Kim, In: Proc. Conf. on Dispersion Strengthened Aluminium Alloys, Y.M. Kim and W.M. Griffith (eds), The Minerals, Metals and Materials Society, Warrendale, PA (1988), p.157. S.M.L. Sastry et al, Met. Trans., A15, 1451 (1984). D.G. Morris and M.A. Morris, In: Prof. Conf. on New Materials by Mechanical Alloying, E. Artz and L. Schultz (eds), Deutsche Geselschaft für Metallkunde, (1989), p.3. C.R. Clark et al, In: Proc. Int. Conf. PM 2 Tech ’95, M.A. Phillips and J. Porter (eds), MPIF, Princeton (1995). Z.A. Zhang et al, Met. and Mater. Trans., 25A, 73 (1994). R. Sundresan and F.H. Froes, Key Eng. Mater., 29-31, 199 (1989). R. Sundresan and F.H. Froes, J. Metals, 39 (8), 22 (1987). E. Hellstern et al, J. Less-common Metals, 140, 1 (1988). J. Eckert and L. Schultz, Mat. Sci. and Engn., A134 1389 (1991). A. Inoue et al, Mat. Trans. JIM, 32 (2) 148 (1990). W.E. Frazier and J. Cook, In: Solid State Powder Processing, A.H. Clauer and J.J. deBarbadillo (eds), The Minerals, Metals and Materials Society, Warrendale, PA (1990), p.257. M.L. Ovecoglu and W.D. Nix, In: New Materials by Mechanical Alloying Techniques, E. Artz and L. Schultz (eds), Deutsche Gesellschaft fur Metalkunde, Oberusel, Germany (1989), p.287. M.L. Ovecoglu et al, Met. and Mat. Trans., 27A, 1033 (1996). P.S. Gilman and K.K. Sankaran, Ibid Ref.5. R. Sundares and and F.H. Froes, 6th World Conference on Titanium, Cannes, France (1988).

Questions 1. In what way are the two techniques MA and RS similar? 2. In what do mechanically alloyed materials differ from rapidly solidified ones? 3. Give two examples of amorphization reaction which can be achieved by MA but not by RS. 4. Explain how magnesium can be added to titanium? 5. How Al–Li alloys produced by RS techniques from from that of MA. 6. What advantages are associated with MA of rapidly solidified materials? Explain the phenomenon with the help of suitable examples.

102

Applications

12 APPLICATIONS Since its advent, the MA technique has been applied to develop various novel compositions and improve the performance of existing materials. Many of these are produced at industrial level and find applications in the commercial sector, while many are potential candidates to find an application. This is discussed in the following section. 12.1 NICKEL-BASE SUPERALLOYS Four nickel-base superalloys are now produced commercially using the MA technique [1]. These are INCONEL MA 754, MA 758, MA 6000 and MA 760. The nominal composition of these alloys is given in Table 12.1. MA 754 has a solid solution strengthened Ni–20% Cr matrix with only low levels of titanium (0.5%) and aluminium (0.3%) along with 0.6% Y 2O 3 as the dispersoid. It can withstand up to 1100°C. MA 758 has been developed for resistance to molten glass corrosion. In MA 6000 alloy, tungsten and molybdenum provide solid solution strengthening. Chromium and aluminium with titanium, tantalum and tungsten improve oxidation and sulphidation resistance. MA 760 is a relatively new variant of commercial alloys prepared by MA. It has been designed for the purpose of achieving a balance of high-temperature strength, long-term microstability and oxidation resistance. Batches of mechanically alloyed superalloy powders are produced in a commercial size attritor mill under controlled conditions. A mixture of powders of nickel (particle size = 5 µm), chromium, molybdenum, tungsten, tantalum and a master nickel-base alloy powder containing the reactive elements aluminium, titanium, boron and zirconium Table 12.1 Composition of MA nickel-base superalloys* Alloy

Ni

Cr

Al

Ti

Mo

W

Y 2O 3

MA 754 MA 758 MA 6000 MA 760

Bal. Bal. Bal. Bal.

20 30 15 20

0.3 0.3 4.5 6.0

0.5 0.5 2.5 –

– – 2.0 2.0

– – 4.0 3.5

0.6 0.6 1.1 0.95

* Products of Inco Alloys International Inc

103

Mechanical Alloying

Fig.12.1 Basic process of consolidation and recrystallization for MA nickel superalloys.

(particle size 150 µm) is used. Yttrium oxide is introduced into the mixture in the form of 1 µm aggregates, each consisting of numerous particles of 20 to 40 nm diameter. The MA powder is then degassed in vacuum at 538°C and sealed in mild steel cans. These cans are then extruded using an extrusion press at temperatures in the range 1010– 1066°C and at an extrusion ratio of 13:1 (Fig.12.1). The microstructure at this stage consists of incredibly fine (0.4 µm) equiaxed grains resulting from dynamic recrystallization during the hot deformation. Hot rolling is carried out at 1025°C with 20% thickness reductions per pass with a total thickness reduction ranging up to 90% [2]. Afterwards, zone annealing is carried out at 1230–1290°C, which increases the grain aspect ratio to more than 15 by a process of secondary recrystallization, giving a structure reminiscent of directional solidification [3] and hence the high-temperature strength of the alloy. In general, 40–60% thickness reductions are required to produce excellent directional recrystallization. The recrystallized grains tend to assume plate-shaped morphologies (see Fig.7.4).

Fig.12.2 TEM micrograph of recrystallized MA 760 superalloy, cooling rate 10°C per min (courtesy Dr H.K.D.H. Bhadeshia). 104

Applications

Microstructural observation shows that γ′ precipitates when fully recrystallized appear as ‘cubes’ in TEM (Fig.12.2). In general, three types of particles Y 2 O 3, Al 2O 3–Y 2O 3 garnets and chromium carbides form during recrystallization rather than during extrusion [4]. Chromium carbides are found to be present in large quantities in the extruded material. These carbides get dissolved during heating to recrystallization and can reprecipitate during subsequent cooling if the alloy sample is not recrystallized. In recrystallized samples, carbides do not reprecipitate rapidly because of the dearth of grain boundary nucleation sites. Inhomogeneous distribution of particles appears in forms of stringers and also as carbide depletion zones (Fig.7.4). However, particle depletion has a small influence on the local microhardness. The properties of these MA superalloys are shown in Table 12.2. They possess excellent high- and low-cycle fatigue resistance strength (~1300 MPa) and ductility (~3%). Comparison of 1000 hrs specific rupture strength with three established directionally solidified and single crystal alloys (Fig.12.3) shows that these materials are superior above 900°C. The failure mechanism at elevated temperature is intergranular fracture along transverse grain boundaries nucleated by cavities that form during grain boundary sliding. Nucleation of voids is retarded due to diffusional accommodation of grain boundary sliding. Depletion of chromium, aluminium and titanium in the surface zones, due to prefTable 12.2 Typical properties of MA nickel-base superalloys (Ref.5) Alloy Property MA 754

MA 758

MA 6000

MA 760

Ni–Cr

Ni–Cr

Ni–Cr– γ '

Ni–Cr–γ '

8.3

8.14

8.11

7.88

Young's modulus at 20°C, GPa

149

–

203

–

Yield strength, 0.2% offset, MPa (at 1095°C)

134

147

192

140

Tensile strength, MPa (at 1095°C)

148

153

222

141

Elongation, % (at 1095°C)

12.5

9

9

15

Alloy type Density, g/cm

3

Stress to rupture, MPa (at 1095°C) 100 hr

102

50

131

115

1000 hr

94

–

127

105

105

Mechanical Alloying

Fig.12.3 Stress for 1000 hours life to rupture as a function of temperature for alloys MA 6000, and other reference material.

Fig.12.4 (right) Comparison of oxidation and sulphidation resistance of MA 6000 with other superalloys (Ref.2).

erential evaporation, also contributes to initiation of the failure. Figure 12.4 shows that MA 6000 alloy possesses excellent oxidation resistance and sulphidation resistance equal to IN-100 and IN-792, respectively, which can be attributed to the homogeneous distribution of the alloying elements and the improved scale adherence due to the dispersoid itself. The mechanical properties of the MA materials combined with their resistance to high-temperature corrosion/oxidation attack offer a level of performance that cannot be obtained in conventional alloys. 106

Applications

The MA nickel-base superalloys are difficult to form. However, good progress has been made in forging and forged components which are in commercial use. Other forming operations such as hot shear forming, ring rolling and hot upsetting have been done at the stage when alloy has low strength and high ductility in fine grain conditions. The MA 754 alloy has been used as components in the hot section of military jet engines and gas turbine vanes at operating temperatures above 1000°C (Fig.12.5). Many of the recently developed single crystal alloys are unable to perform in these type of applications unless they are plasma sprayed with MCrAl(Y) type materials to improve the oxidation resistance. The MA 760 alloy is mainly for the use in gas turbine components. The alloy combines excellent hot corrosion resistance with long term, high-temperature strength that is beyond the capabilities of conventional superalloys.

Fig.12.5 Brazed aircraf, gas-turbine vane assembly fabrication of MA 754 superalloy.

12.2 MA STEELS Many MA materials are now commercially available, but steels produced using MA show particular promise in a variety of applications. For example, MA 956 (Fe–20% Cr–4.5% Al–0.5% Ti–0.5% Y2O3) is a chromium-rich, ferritic stainless steel containing aluminium for oxidation resistance, (aluminium forms an Al 2O 3 protective layer at high temperatures) together with a dispersion of Y 2O 3 particles for creep resistance. An alternative ferritic stainless steel variant MA 957 (Fe–13.5%Cr– 0.3%Mo–1%Ti–0.3%Y2O3) contains titanium rather than aluminium, and is designed for application in the nuclear industry. MA 956 can be used at operating temperatures of over 1300°C in corrosive atmospheres. The 107

Mechanical Alloying

a

b Fig.12.6 Components made of MA 956: a) burner nozzles machined from forgings; b) air-stream swirlers fabricated from sheet.

Fig.12.7 Use of MA 956 in heat treatment equipment: a) vacuum furnace hearth; b) heat treatment furnace basket fabricated from rods.

alloy is used as sheet material for aircraft and industrial gas turbine combustors, in burners, swirlers and heat exchangers of power generation equipment (Fig.12.6), and in heat treatment equipment (Fig.12.7) [5,1]. 108

Applications

INCONEL MA 957 alloy has been developed for an application as nuclear fuel cladding material in a fast breeder reactor. Conventional austenitic alloys can not be used in this application due to swelling caused by the high neutron fluxes, while conventional ferritic steels generally have inadequate creep strength at the service temperature of 700°C. These steels are produced basically to meet these requirements. The steels after MA and consolidation have an ultrafine microstructure containing submicron grains of ferrite. The hardness in this condition is unacceptably high (in contrast to the MA nickel-base superalloys), so the steels are usually recrystallized into a coarse directionally recrystallized grain structure which is also ideal for elevated temperature applications where creep resistance is of prime importance. The MA 956 steel exhibits a strong {110} fibre texture in the extruded condition, while MA 957 has a relatively poor texture. However, after recrystallization there was a significant difference in texture which cannot be explained on the basis of deformation or recrystallization theory. Chou and Bhadeshia [6] revealed that the ferrite present in MA 957 partially transforms to austenite at the extrusion temperature (970 to 1010°C) because of its relatively low chromium concentration and the absence of aluminium. The austenite then transforms to martensite on cooling, giving a new set of orientations and, consequently, a more random crystallographic texture. During annealing, recrystallization nucleates via both a subgrain coalescence mechanism and grain-boundary migration. The existence of austenite phase enhances the development of cubic texture. Dour Metal have developed a series of oxide dispersion microforged material (ODM) iron base FeCrAl alloys, by MA, which meet the requirements of strength and oxidation resistance at high temperatures [7]. The composition of these alloys is given in Table 12.3, while a comparison of their properties is given in Table 12.4. These materials have good hot forming properties in the microcrystalline condition as compared to conventional FeCrAl alloys. The ODM has been found suitTable 12.3 Composition of Fe–x%Cr–y%Al–1.5%Mg–0.6%Ti–0.5%Y2 O 3 alloys Alloy

Cr%

Al%

ODM 331

13

3

ODM 361

13

6

ODM 031

20

3

ODM 061

20

6

ODM 751

16.5

4.5

109

Mechanical Alloying Table 12.4 Hot tensile properties of ODM alloys Microcrystal 900°C

Macrograin 900°C

Alloy

YS, 0.2% Offset (MPa)

UTS (MPa)

Elongation (%)

YS, 0.2% Offset (MPa)

UTS (MPa)

Elongation (%)

331

28

41

157

169

176

10

361

42

61

86

166

177

12

061

49

71

88

159

173

16

751

32

47

134

–

–

–

FeCrAl

–

35

–

–

–

–

MA956

–

–

–

108

115

8

F10R

–

–

–

–

165

16

able even at temperatures above 1100°C, which is the upper limit for the use of superalloys for both the mechanical properties and the oxidation resistance. A suitable heat treatment of ODM has also been developed to favour the transgranular rupture in these materials, which improves ductility. When compared with MA 956, the time to rupture is multiplied by a factor of 100. The superior mechanical properties of these ODMs above 1000°C is due to the strengthening effect of dispersoids. Thus, it is envisaged that these ODMs may find applications in energy conversion systems (particularly of heat exchanger tubes) as well as other high temperature applications in corrosive atmospheres. For the requirements of modern combustion engines, high wear resistance steel having a tough matrix and hard, fine dispersoids is required. In this respect, four hard phases can be used, namely: NbC, TiC, TiN and Al 2O 3 . All other hard phases are either too soft or unstable within the steel matrix at sintering temperatures. Sintered steels containing hard phase in the range of (10–25 vol.%) have been produced by MA of elemental composition in an attritor mill followed by liquid phase sintering at 1280–1320°C in a hydrogen atmosphere. A high phosphorous content leads to a harmful embrittlement of the grain boundaries, as a consequence the phosphorous content is chosen very carefully [8,9]. The influence of this hard phase on the mechanical properties of the composite material is given in Table 12.5. It can be observed that the effect of wettabilities of different hard phases with the matrix on sintered densities of the composite material is much less. However, stiffness of the composite increases as a function of Young’s modulus of the reinforcing component. The strength of Al2O3-containing 110

Applications Table 12.5 Properties of the used hard phases and the composite materials with a composition of Fe–10 vol% hard phase – 0.6% P–0.9% C NbC

TiC

TiN

Al2O3

Hard phase density (g/cm 3)

7.78

4.98

5.40

3.98

Young's modulus (kN/mm 2)

580

470

590

410

Composite material theoretical density (g/cm 3)

7.85

7.57

7.61

7.47

Sintered density achieved (% TD)

97.9

98.2

97.4

97.46

1008+88

995+66

806+64

742+21

Elongation (%)

3.7

3.2

1.8

1.6

Young's modulus (kN/mm 2)

229

218

205

215

Tensile strength (N/mm 2)

material is 25% less compared to that of NbC, and the elongation fracture is 50% less. Surprisingly, sintering densities are found to be increased with the amount of an inert second phase in the composite which does not take part in the sintering process. It has been attributed to the fact that sintering in the composite materials takes place by means of intergranular diffusion for a longer period of time while the grain size remains more or less stable. The wear behaviour of these materials has been found to be the best when the hard phase content is about 10 vol.%. It is also reported that dry wear is reduced using coarser hard phase. For oil lubrication at higher surface pressures (800 N/mm 2) and longer times (100 hrs), finer hard phase particles lead to less wear. Thus, for oil lubricated tribosystems, these hard phase containing sintered steels offer a better alternative to conventional wear resistance PM steels like sintered and forged T15 MSS. Moreover, the matrix hardness can be controlled within a wide range by heat treatment and, hence, can be matched to the respective application. 12.3 ALUMINIUM-BASE MATERIALS The MA process can be applied to aluminium-base systems provided that organic process control agents are used to control the extreme tendency of aluminium to weld to itself during high energy milling. All MA aluminium products have highly uniform dispersion of equiaxed oxide (Al2O 3) and carbide (Al 4C 3) particles, formed due to the reaction of aluminium with oxygen and carbon (formed as residue of PCAs) particles during the consolidation. These dispersoids of extremely fine sized oxide and carbide particles (30–40 nm) stabilise a very fine equiaxed grain (0.2 to 0.5 nm) and dislocation substructure during thermomech111

Mechanical Alloying

Fig.12.8 Flow sheet for processing of MA aluminium alloys (Ref.11).

anical processing [10]. In addition to the strengthening caused by these finely dispersed oxides and carbides, a significant part of the strengthening in these alloys is due to the fine grain size and high dislocation density resulting from the severe working of the powders. Further benefits can be obtained in these alloys by adding small amounts of carbideforming elements such as titanium. A typical flow sheet for processing MA aluminium alloys is shown in Fig.12.8 [11]. The development of MA aluminium alloys can be put into three classes: high strength alloys, low density alloys and high temperature alloys. High strength alloys The major application of MA aluminium alloys is at ambient temperatures rather than elevated temperatures. One example is IN 9052 (Al–4% Mg–1.1% C–0.8% O) (Fig.12.9). Although a solid solution alloy, IN 9052 gives a combination of tensile strength, fatigue strength and toughness normally seen only in precipitation hardened 7000 series alloys. The high strength of this alloy arises from magnesium solid solution strengthening, ultrafine grain size, finely dispersed MgO, Al2O3 and Al 4 C 3 and substructural strengthening [12]. Strengthening due to precipitation can be superimposed on the above mechanisms when higher strength is essential but corrosion resistance is of secondary importance. Because it is a solid solution material, IN 9052 also possesses 112

Applications

Fig.12.9 TEM micrograph showing fine-grain structure of IN 9052.

Fig.12.10 Strength and corrosion resistance of IN 9052 compared with conventionally made alloys (Ref.11).

excellent resistance to corrosion and stress corrosion cracking (Fig.12.10). The fatigue crack growth rates of IN 9052 and another mechanically alloyed aluminium alloy IN 9021-T4 (Al–4%Cu–1.5%Mg– 1.1%C–0.8%O) are comparable with that of 7075 alloy. These alloys are being considered suitable for marine and metal matrix composites. Low-density alloys For the aerospace industry, an alloy with properties equivalent to or better than those of the widely used 7075-T73 alloy and with lower density is desired. For density reduction in aluminium, there are generally only two choices of alloying additions that merit practical consideration – lithium and magnesium. The proportion of these elements which can be tolerated is limited by the possible occurrence of stress-corrosion 113

Mechanical Alloying

cracking at levels above the solubility limits. The possibility of achieving up to 15% reduction in density along with up to 10% increase in elastic modulus (amounting to a combined increase of approximately 30% in specific strength) by alloying aluminium with lithium is possible. The aluminium–lithium system seems to hold high promise in the aerospace industry and because of its ‘plasma-confinement’ properties, it also has prospects in nuclear engineering [13]. The existing production methods of casting and powder-making by atomisation lead to a material of inhomogeneous composition structure. Lithium or lithium compounds (Al 3Li) are segregated at the grain boundaries, causing excessive brittleness of the material and reducing its resistance to fatigue. The MA process has specific advantages in the case of highly reactive constituents (like Al, Mg and Li) under atmospheric conditions as well as for constituents which have high partial vapour pressure at their liquid temperatures, as is the case with lithium. The difficulties inherent in the production process are due to the wide disparity between the three constituent metals. Pure aluminium is ductile, while magnesium is relatively brittle. Fine magnesium powder is pyrophoric and tends to ignite spontaneously under ambient conditions. Pure lithium metal cannot be handled directly, and must always be kept under a protective liquid (e.g. paraffin oil). It is very soft and has a high rate of sublimation at moderately high temperatures. Al–3% Li–2.2% Mg has been produced by carrying out MA in a vibratory mill in a vacuum of 10 –5 torr. The master alloy powder is first produced by milling a mixture of Mg powder and Al–Li powder. The aluminium powder is used as a diluent. The difficulty presented by the tendency of the soft aluminium particles to agglomerate and adhere to the grinding medium of the ball-mill is removed by the presence of a brittle (both magnesium and AlLi) constituents, which interposes itself between the aluminium particles and keeps them apart [14]. Flow-sheet of the process is shown in Fig.12.11. The X-ray diffraction (Fig.12.12.) showed that Al–Li undergoes partial solution after grinding and full solution following hot pressing. The Mg is found to distribute uniformly in the alloy. In the fine grains, Li diffuses to the grain boundaries forming a β−(Al–Li) phase which causes brittleness because of its incoherence. However, in the case of larger grains, coherent precipitates of δ′-(Al 3Li) are coherent with the matrix. Thus, MA provides a microstructure with coarse grains having a precipitate of hard δ′-(Al3Li) and very fine grains, which improves the mechanical properties at room temperature. After solution treatment at 450°C and ageing at 150°C, the alloy achieves the highest value of specific modulus of elasticity E/ρ (≈28). This makes the alloy suitable for aviation application which can be ascribed mainly to the coherent precipitation of δ′−(Al3 Li) and to 114

Applications

Fig.12.11 Flow sheet of the production process for Al–Li–Mg alloy (Ref.14, Powd. Met. Int., 22(3), 22 (1990) (Verlag Schmid)).

Fig.12.12 X-ray diffraction pattern of an Al–2% Li–4.4% Mg powder: a) before MA; b) after MA; c) after hot pressing (Ref.14). 115

Mechanical Alloying

the low density of the lithium, the lightest existing metallic element. Based on the work of S.J. Domachie et al [15,16], IncoMAP has developed the IN-905 XL alloy (Al–4% Mg–1.5% Li–1.2% C–0.4%O) which matches 7075-T73 in strength but is 8% less dense and 15% more stiff. The outstanding feature of the alloy is the minimal degradation of properties when stressed in the transverse direction. This characteristic is of great importance in forgings which are uniaxially stressed. Its general corrosion resistance is about 100 fold better than that of 7075-T73. Heating cycles up to 450°C do not alter the strength. This characteristic can be important in service as well in manufacturing. Figure 12.13 shows a prototype aircraft wheel forged of this alloy.

Fig.12.13 Prototype aircraft wheel forged of IN 905 XL.

High-temperature alloys The development of MA high-temperature aluminium alloy can be put in three classes: – DISPAL – Al–Ti alloys – Al–Fe–Ce alloys The MA process has lead to the development of Dispersion Strengthened Aluminium (DISPAL) high-temperature aluminium alloys [17]. The alloy powders are processed directly by reaction milling with carbon black and controlled additions of oxygen from the milling atmosphere and with this independent control of the dispersoid content, the alloys can suitably be tailored for applications. Already, forged automotive pistons, high-temperature electrical conductors, interferometers and structural parts in aerospace cryogenic fields have been made from various DISPAL alloys. 116

Applications

To replace titanium in aerospace substructures and the cooler parts of advanced jet engines, Al–(7.5–15)% Ti alloys, e.g. E5(Al–11.6% Ti– 1.9% C–0.7% O), have been developed. These alloys consist of a matrix of aluminium containing a second phase of Al3Ti particles. Several characteristics of Al–Ti alloys make them potentially attractive for engineering use. The high melting point of Al 3Ti and the low solubility and low diffusion rate of titanium in aluminium make them suitable for elevated temperatures. In addition, increased stiffness can be derived from the high elastic modulus of Al 3Ti. When adding more than 0.15% (0.15% amount is sufficient for grain refinement), Al 3Ti particles precipitate directly from the melt and grow to large sizes, which decreases ductility. Even rapid solidification techniques have failed to control the phenomenon as the addition of Ti greatly increases the liquidus temperature. The MA of elemental aluminium and titanium powders distributes titanium uniformly in the matrix of aluminium. Alloying is carried out in a chilled water-cooled attritor mill in a flowing argon atmosphere. The MA powder is degassed at 550°C and extruded into bars at the same temperature with an extrusion ratio of 25:1. During heating for consolidation, the titanium and aluminium react in situ to form Al3Ti particles about 20–250 nm in diameter, providing a structure not possible with conventional manufacturing methods. A MA Al–12.5%Ti (E6) alloy has an elastic modulus of 50% greater than that of conventional aluminium alloys, accompanied by a density increase of 4% with good ductility and high strength at elevated temperatures [18]. The high strength of the MA-processed alloy (grain size 100–300 nm) is obtained by dispersion strengthening through the Al 2O 3 (30–100 nm), Al4C 3 (30100 nm) and Al3 Ti (20–250 nm), by the addition of alloying elements and due to work hardening [12]. It is reported that Al 3Ti forms above 260°C during degassing. These precipitates are stable thermally up to 1000 hrs at 300°C and 500 hrs at 400°C [19]. The tensile strength of Al–10%Ti (E4) alloy has been reported to be almost the same as the conventional precipitation-hardened aluminium alloys (2014-T6, 2219T18) up to 100°C, but it is three times higher above 200°C (Fig.12.14) due to the presence of the dispersoids. The stress rupture strength of the alloy at 300°C is three times as much as the conventional precipitation hardened alloy. The activation energy for creep is 228.34 kJ/mol, indicating the strong interaction between the precipitates and the dislocations. The fatigue strength of MA Al–10% Ti at room temperature is somewhat lower than that of 7075 which also has high tensile strength, but higher than that of 2024. The fatigue strength is found to decrease with an increase in temperature, but much slower than the decrease of tensile strength. In elevated temperature use (472°C), MA alloys are most promising as is a combination 117

Mechanical Alloying

Fig.12.14 Strength at elevated temperatures of mechanically alloyed Al–10% Ti (Ref.11).

of RS with MA of alloys such as Al–Fe–Ce (as discussed in Section 11.2). At the same time, ductility levels are significantly higher than Al–SiC p composites (see Section 12.12). The good ductility is attributed to good binding between in situ formed Al3 Ti and the matrix in situ contrast to the artificial incorporation in metal-matrix composites (MMCs). Another high temperature alloy Al–8% Fe–4% Ce is prepared by MA using stearic acid as the PCA, degassing at 425°C and finally hot extrusion at 425°C using an extrusion ratio of 6:1 [20]. The MA process introduces a fine-scale uniform dispersion of oxides and carbides. The MA material microstructure was stable even after exposure for 1 hr at 610°C. A comparison of the creep response of the material with the RS one, under the same conditions of stress and temperature, shows (Fig.12.15) that MA material is superior. However, ductility values are low due to cleavage type of fracture. A summary of the composition and mechanical properties of MA aluminium alloys developed is given in Tables 12.6 and 12.7, respectively. For other MA aluminium materials, refer to Sections 12.7, 12.10, 12.11 and 12.12.

Table 12.6 Composition of MA aluminium-base materials Alloy

Al

Mg

Cu

Li

C

O

MA 9021*

Bal.

4

–

–

1.1

0.8

MA 9052*

Bal.

1.5

4

–

1.1

0.8

MA 905XL*

Bal.

4

–

1.5

1.2

0.4

DISPAL* *

Bal.

–

–

–

2.0

2.5

* Product of Inco Alloys International Inc. **Product Krebsöge (Germany) (Germany) ** Product of of Sintermetall Sintermetall Werke, work, Krebsöge

118

Applications Table 12.7 Properties of MA aluminium alloys and other reference materials YS (MPa)

UTS (MPa)

El.(%)

Fracture Toughness (MPa/√m)

Elastic Modulus (GPa)`

Density (gm/cm 3)

Corrosion Rate (mm/year)

Fatigue Strength (MPa)

542 613 326 279 462 503

577 634 542 464 525 572

8 12 24 22 13 10

44 – – 32 30 28.6

76 – – – 80 69

2.68 – – – 2.58 2.80

0.009 – – – 0.003 –

>380 – – – 345 –

634 595

717 645

9 3.5

– –

– –

– –

– –

– –

IM(Al–5%Mg–1.5%Li) Dispal–2

193 340

345 370

18 10

– –

16 80

2.55 2.70

– –

– 110

Dispal–1 (Al-12%Si)

340

380

8

–

85

2.66

–

150

Dispal (Al–12%Si–Cu–Mg–Ni)

340

370

3

–

81

2.70

–

120

Alloy

IN 9052 IN 9021 RS 2024 IN2024 IN 905 XL IN 7075–T73 RS7075 (+1%Ni+1%Fe) MA 7075

Fig.12.15 Comparison of creep rate for non-MA (RS) and MA Al–8% Fe–4% Ce alloy (Ref.20).

12.4 COPPER-BASE MATERIALS Heat-resistant copper alloys having high electrical and/or thermal conductivities are required in a variety of applications, including resistance welding electrode contacts, electrical switches, microwave and X-ray tube components, incandescent lamp lead wires and neutron irradiation targets [21]. The wrought Cu–Cr and Cu–Zr alloys commonly used for elevated temperature applications are strengthened through thermomechanical treatments, which cause precipitation and stabilization of a strain-hardened structure. In such alloys, phenomena such as recovery and recrystallization start taking place at a temperature of about 450°C. Using 119

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the MA technique, it has been possible to extend the amount of chromium in the copper matrix. The use of PCAs is avoided in order to have high electrical and thermal conductivity in the final material by minimising impurities. Cu–5% Cr powder could be prepared by MA of elemental powders. The MA powder was consolidated by hot extrusion. A strength level of 75 MPa could be achieved, which is attributed to fine grain structure and dispersion strengthening. Ductility of 8% and conductivity of 60% of the copper were found after the heat treatment [22,23]. Alumina dispersion-strengthened copper has been produced by MA of copper powder and an amorphous Al(OH) 3 powder [24]. The Al2O 3 is created in situ in the copper matrix at 254°C during heating of the green compacts to the sintering temperatures. Consolidation of the MA powder was carried out by cold compaction followed by sintering at 940°C in a nitrogen atmosphere. The highest hardness (110 VHN) was obtained for the optimized value of 0.65% of Al 2O3 in the matrix showing a significant improvement in comparison with the copper matrix, hardness (48 VHN). It was also interesting to note that use of a crystalline powder as a dispersoid results in hardness lower than in the case of amorphous Al2O 3. The hardness variation of the MA matrix with exposure to high temperature at 940°C for different periods of time is shown in Fig.12.16. With an initial minor drop in hardness due to the recovery of the cold worked state, the material is found capable of retaining its hardness at high temperature exposure. Thus, ODM is capable of retaining hardness/strength up to 0.9 of the absolute melting temperature of the copper. This approach can be used to develop dispersion-strengthened, high-temperature copper alloys. An elemental alloy (Cu–8% Ti–4% B) has been mechanically alloyed [25]. The solubility of Ti in Cu is reported to be 2 at.%. TiB2 is formed

Fig.12.16 Effect of annealing time on softening behaviour of MA Cu–0.98% Al2 O 3, annealing temperature 940°C (Ref.24, Powd. Met. Int., 22 (5), 15 (1990) (Verlag Schmid)). 120

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in the ternary mixture. It is reported that refinement is accelerated by TiB 2 or B. The MA powder was consolidated by sintering at temperatures between 400 to 900°C. The hardness was found to increase with sintering time. High temperature sintering enhanced the densification and formation of TiB 2 but reduced the hardness. The MA process has also resulted in copper-rich solid solutions in the prealloyed Cu–Pb [26], Cu–C [27] and Cu–Al–C [27] systems. For other copper-base materials, refer to Section 12.11. 12.5 TITANIUM SYSTEM The MA titanium base alloys are still in the research state, but some interesting advances have been made in the following: – Ti–Mg immiscible; – dispersion strengthened TixAl and Ti5 Si3 . Though Ti–Mg forms an immiscible system and solubility obtained of Mg in titanium by conventional methods is negligible, it is an attractive alloying element to titanium due to its low density. Using MA, solid solution of Mg up to 3.1% in titanium has been achieved [28]. Magnesium also appears to be a very strong β stabiliser, which depresses the β transus in binary Ti–Mg as well as in ternary Ti–Mg–Gd systems. Titanium aluminides based on γ-TiAl (L1 structure) are potentially very attractive low-density materials for use at elevated temperatures in aerospace applications. However, these alloys, like most intermetallics suffer from low ductility and fracture toughness at ambient temperatures due to the presence of a superlattice, and low corrosion resistance even at the moderately high temperature. However, the inherent brittleness of these alloys limits their fabricality. Alloying additions such as Mn for ductility and Cr and Nb for oxidation resistance improve these properties appreciably. The nanocrystalline equiaxed γ-grains and the fine lamella γ/α 2 microstructure also results in the same effects. The MG process has been applied to two such alloys [29,30], Ti–25% Al– 10% Nb-3% V–2% Mo(at%) and Ti–48% Al–2% Mn–2% Nb(at%). The microstructure of the consolidated gas atomised powder shows a mainly uniform two-phase structure consisting of equiaxed α 2 with a grain size of about 4 µm, with intergranular B2 and some occasional patches of lenticular α 2 grains. However, MG of powder prior to consolidation results in a structure consisting mainly of very fine grains of about 1 µm, with some coarser grains of about 3 µm retained (Fig.12.17). Dispersoids in the consolidated alloy are coarse, about 0.1 µm in size. An appreciable improvement in the hardness of the MA material (720 KHN) is noticed as compared to atomised powder (340 KHN), which is attributed to the refinement of the microstructure and the transfor121

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Fig.12.17 Optical micrographs of compacts of MA Ti–25% Al–10% Nb–3% V–2% Mo: a) fine grained structure; b) coarser grains, containing a larger portion of lenticular grains (Ref.29, Metal Powder Report, 44 (3), 206 (1989) (Elsevier Science)).

mation of α2 phase to B2 phase. The microstructure of atomised powder has completely B2 structure, which turns to a mixture of both α 2 and B2 phases on consolidation. The proportion of B2, a more ductile phase, has been found to be more in the case of MA powder as compared to atomised powder. The Ti5Si 3 intermetallic compound is also a promising high temperature material for aerospace application. However, the inherent brittleness of this alloy limits its formability. The MA process has been found to be suitable to produce this alloy in an amorphous and/or nanocrystalline form to improve its mechanical properties [31]. The MA of elemental powders was carried out in a Spex 8000 mixer mill in argon atmosphere using hexane as a PCA medium, until a completely amorphised phase was produced (25 hrs). The DTA results show that the powder has a crystallisation temperature of 640°C. The MA powder is then consolidated using a modified shock consolidated method. The modification includes a cylindrical gap provided around the powder sample region of the capsules to prevent radial focusing of shockwaves to minimize generation of very high and non-uniform pressures within the axial regions of the sample. Thus, the average calculated peak pressure in the powder mixtures packed to 64% theoretical density, for an impact velocity of 0.9 km/sec, is in the range of 20 to 40 GPa, with the onedimensional loading pressure being 5 GPa. During compaction, a significant amount of crystallization, forming 30 to 40 nm crystals of 122

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orthorhombic TiSi 2 and Ti 5Si 3 intermetallic compounds, has been reported. This partial crystallisation of the amorphous powder during shock compact occurs due to mean-bulk shock temperatures exceeding 1000°C (recrystallization temperature = 640°C). The compacts are subsequently annealed above the crystallization temperature at 800°C for 1 hr, which results in limited growth to 50 nm crystallite size. The microhardness of the shock-compact was 1100 KHN, which increases to 1250 KHN upon subsequent annealing with the formation of a more homogeneous nanocrystalline microstructure. Nanocrystalline Ti–N powder has also been prepared by reaction milling in a nitrogen atmosphere. The formation of TiN was found to depend on Ti–N solid solution. A strong solid solution hardening has been detected [32] both in the as-milled and the annealed (at 350 to 540°C) nanocrystalline Ti–N alloys. The approach may probably be used in improving the titanium aluminides. Two advanced β alloys Ti–1% Al–8% V–5% Fe–1% Er and Ti–24% V–10% Cr–5% Er, both with extensive rare earth additions, are being developed for subsequent oxide dispersions by a combination of RS and MA, as discussed in Section 11.2. 12.6 MAGNESIUM-BASE MATERIALS The MA process has been used to alloy magnesium with Fe/Cu/Cr/C/ Ti for developing ‘supercorroding’ alloys, which are used for making links with precisely controllable corrosion rates for releasing deep sea equipment at specific depths [17]. Such supercorroding alloys have also been of use in other submarine applications, such as a heat source in diver suits and as a hydrogen gas generator. The advantage of MA in producing a supercorroding alloy is the extent of control afforded by the closeness of the anode/cathode metals achieved. By adjusting the alloy composition, the reaction rates can be precisely controlled in the device (Fig.12.18). Mg–M is a promising system for hydrogen storage. One of the ways to kinetically improve magnesium-based hydrogen storage material is by the addition of metals and alloys whose hydrides have a high dissociation tension. By alloying with catalyst elements/intermetallic compounds, the hydrogenation capability of magnesium is substantially enhanced [17]. The additions include elements such as: − Ni, Ce, La which can be alloyed with Mg conventionally. These hydride-forming elements can function as 'hydrogen pumps' owing to stoichiometric variation. − Elements such as Fe, Co, Ti, Nb that do not normally alloy but form intermetallics with magnesium which can absorb and deabsorb hydrogen. 123

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Fig.12.18 Corrosion rates in 'supercorroding' MA magnesium alloys (Ref.17, Metal Powder Report, 44(3), 195 (1989) (Elsevier Science)).

− Intermetallics such as LaNi5, Mg2Cu, CeMg12 and even oxides such as TiO 2. These are generally prepared by argon melting. The MA technique has been used to produce the first and second type materials with improved performance (Fig.12.19). The MA is carried out in a planetary ball mill in an argon or hydrogen atmosphere. Severe cold work resulting from MA aids diffusion by providing many sites with low activation energy of diffusion. In addition, the close mixture of the powder constituents decreases the diffusion distances to the micrometer range. Precipitation may occur or metastable phase may form throughout the powder particles, or the dispersoids may provide nucleation sites for the formation of hydrides. For example, the MA of an elemental magnesium and nickel alloy results in dispersion of the MgO film present on the surface of magnesium particles, and disordering of the surface of magnesium and nickel particles in the matrix. The XPS results [33] show that the nickel atoms present in the subsurface of MA powder

Fig.12.19 Kinetics of five-cycle hydriding and dehydriding: 1) MA Mg; 2) MA (Mg5% Nb); 3) MA (Mg–5% Fe)/Mg–5% Co)/(Mg–5% Ni); 4) MA (Mg–5% Ti); 5) MG (Mg 2 Ni); 6) Mg (P=0.7 MPa, T=85°C) (20 µm); 7) MA (Mg–50% C); 8) Mg–25% Ni (conventional) (Ref.33). 124

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have a binding energy of Ni 2p 3/2 less by 0.5 eV, which is the ratedetermining step for dissociative adsorption of hydrogen on nickel atoms. The formation of Mg 2Ni is found to take place in the hydriding reaction of the MA alloy, and practically all the nickel converts into the intermetallic compound within 2–3 cycles. Mg 2 Ni is shown to form mainly by 2MgH 2 + Ni → Mg 2Ni + 2H 2 Thus, the presence of these alloying elements in the MA powder surface, hydrogen adsorption and magnesium hydride nucleation are essentially accelerated, and the reaction begins at maximum rate. However, the formation of a ternary hydride like MgFeHx (x = 6) decreases the hydrogenation and the formation of a stable ternary hydride like Mg 2CoH x (4 < x < 5) is not favourable for dehydriding kinetics. 12.7 SUPERSATURATED SOLUTIONS The use of the SSE to produce supersaturated solutions without producing a second phase can be achieved by various techniques like: solid state quenching, rapid solidification, condensation from the vapour state and irradiation/ion implementation. Supersaturated solid solutions are stronger and harder than equilibrium solid solutions, and on decomposition can produce a higher volume fraction of fine second phase particles. The actual values for the SSE are calculated using Vegard’s law, which states that the lattice parameter and the amount of solute added to the solvent lattice have a direct, linear relationship. The technique of MA has been applied for the SSE in a variety of systems as discussed below. A non-equilibrium supersaturated fcc solid solution is formed in an Ag–Cu system in which the entire composition range has positive heat of mixing and is immiscible to each other by MA [34]. The supersaturated solid solution formation in this system means the elevation of free energy caused not only by the stored energy as defects but also by the effect of mixing of immiscible materials in dimensions as fine as that of a nanometer order. Decomposition of the supersaturated solid solution in MA alloy takes place at above 157°C which can be identified by the accompanied sharp decrease in electrical resistivity (Fig.12.20). At the initial stage of decomposition, the grain size of the decomposed phases remains at nanometer size. Upon heating above about 327°C grain growth takes place. The SSE has also been achieved of Mg, Cu and Si in aluminium, as well as of Al in magnesium as discussed in Section 11.1. A supersaturated nanocrystalline Al–8% Ti–0.3% Fe (at.%) alloy powder 125

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Fig.12.20 Changes of the electrical resistivity with temperature in Ag–Cu alloys (Ref.34).

has also been prepared by MA [35]. The equilibrium solid solubilities of V–Cu and V–Fe alloys are low but V–30 at% Cu and V–50 at.% Fe alloys are reported to form solid solution after MA [36]. 12.8 MAGNETIC MATERIALS Nd–Fe–B magnets The advent of high-performance permanent magnets based on Nd–FeB in 1983 has generated a huge scientific and technological interest. The main reason for this trend is the cheapness and abundance of the involved elements. The origin of the best magnetic properties in these magnets is the tetragonal structure of Nd 2 Fe 14B phase. The Curie temperature (T c) of these magnets is 314°C which can be improved by a small addition of cobalt. An anisotropic field H A of about 6000 kA/m enables high coercivities (800 to 200 kA/m) and a saturation magnetization of 5120 mA/m. An MA technique using a planetary mill has been recently used to produce isotropic and anisotropic Nd–Fe–B magnets [37]. The MA produces a closely mixed powder of Nd, Fe and B, which has low coercivity due to the presence of α-Fe magnetic phase which is transformed to hard magnetic Nd 2Fe14B phase by a solid state reaction by an annealing treatment typically at 700°C for 30 min. These isotropic microcrystalline powders are used for making bonded type isotropic and compacted anisotropic magnets (Fig.12.21). Compaction is carried out by hot pressing at 700°C for 5 min under a pressure of 1 kbar. Density levels of 98% of theoretical values could be obtained using the process of die upsetting where a hot pressed sample is uniaxially compressed to half its height in a die whose diameter is √2 times the sample. The crystallographic texture present in the consolidated sample is attributed to stress-induced grain growth and grain rotation during the compression. 126

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Fig.12.21 Preparation of Nd–Fe–B magnets from mechanically alloyed powder (Ref.37).

The ratio of remanences (measured parallel and perpendicular to the press direction) critically depends on the deformation rate and the number of deformation steps. A value of 2.6 could be obtained for this ratio. Thus, flake geometry, as obtained in rapidly quenched material, is not essential for the formation of anisotropic samples, since the mechanically alloyed powder has a highly irregular shape. Figure 12.22 shows that the coercivity of the hot pressed samples is slightly lower, but shows the same composition dependence. With an increase in neodymium content, a higher coercivity of up to 1400 kA/m at x = 9 is obtained with an increase of the second phase after x = 6. Sm–Fe magnets Magnetic alloy systems like Sm (Fe,X)12 and (Sm,Zr)Fe3 have also been prepared successfully using an MA technique [38,39]. The MA of elemental samarium and iron powders leads to a two-phase mixture of amorphous Sm–Fe and α-Fe. Heat treatment at 650–800°C for several hours produces Sm 2 Fe 17, with a Th 2Zn17 crystal structure, which has a Curie temperature of only 116°C and an in-plane anisotropy. In this state, the material is still soft magnetic because of the in-plane anisotropy of Sm2Fe17. During the subsequent nitriding treatment at 400550°C in a N 2 atmosphere, nitrogen atoms are introduced into its Th 2Zn 17 crystal structure, resulting in an extension of the lattice parameter as calculated by the X-ray diffraction patterns (Fig.12.23b and

Fig.12.22 Demagnitization curves of resin-bonded (MM1), compacted isotropic (MM2) and anisostropic (MM3) Nd16Fe76B8 magnets (Ref.37). 127

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12.23c), to form Sm 2Fe 17N x phase [40]. The amount of residual α-Fe depends on the Sm content, and its optimum amount has been found to be 12.5 at.%. The grain size after milling was typically 10–30 nm and increased to 50 nm after annealing. The magnetic domain size was found to be 300 nm. Oxygen was found to enhance precipitation of α-Fe and a maximum of 9000 ppm is suggested [41]. Low oxygen is found to suppress precipitation of α-Fe which deteriorates the magnetic properties. This new phase, Sm 2Fe 17N x, has coercivities up to 2400 kA/ m and uniaxial anisotropy. The phase has also improved the Curie temperature T c (470°C compared to 320°C), the saturation magnetization M s, (1.6 T compared to 1.54 T) and the anisotropy field H A (14 T compared to 8 T) as compared to Nd–Fe–B magnets. Therefore, magnets of Sm–Fe–N should have the potential to compare favourably with wellestablished Nd–Fe–B magnets, if the high anisotropy field can be explained to create a high coercivity and a processing route can be established. An isotropic resin-bonded (Sm 12.5Fe 87.5) 1–xN x magnet shows remanence M r = 0.71, which is the same for mechanically alloyed Nd 16 Fe 76 B 8 . The temperature dependence of the coercivity and remanence of this phase is much smaller than that for Nd-Fe-B. At 150°C, the coercivity is still 1450 kA/m. The major drawback of this new hard magnetic material is its limited stability at elevated temperature. It has been found that 2:17 nitride is stable up to 650°C, thereafter it decomposes into samarium nitride and α-Fe. This instability of the 2:17 nitride prevents the application of sintering technique used for NdFe–B or Sm–Co for the production of anisotropic magnets. Sm–Co magnets The MA process has also been used to produce conventional SmCo 5 and Sm 2 Co 17 magnetic materials. To produce these materials, MA of

Fig.12.23 X-ray diffraction diagrams of Sm 12.5 Fe 87.5 : a) after MA; b) after the formation of Sm2 Fe17 by annealing in vacuum; c) after the nitriding reaction (Ref.39). 128

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Sm 16.7 Co 83.3 and Sm 10.5Co 89.5 powders, corresponding to 1:5 and 2:17 stoichiometry, respectively, is performed in a planetary ball mill under an argon atmosphere [42]. The MA in the case of Sm 16.7Co 83.3 produces an amorphous phase possibly with traces of free Co, and for Sm 10.5 Co 89.5 a two-phase microstructure of an amorphous matrix and Co solid solution. The two-phase region in the case of Sm 10.5Co 89.5 results from a metastable equilibrium between the amorphous phase and the Co solid solution. The crystallization of amorphous state to the 1:5 structure, SmCo 5 , occurs at about 550°C for both alloys, which has a high coercivity of about 2200 kA/m and a remanence of 0.52T. The degradation of the remanence M r results at a high annealing temperature due to Sm losses during annealing accompanied by the precipitation of Co. The coercivity of Sm 2Co 17 is found to be low due to the lower anisotropy field of 2:17 phase. To increase the coercivity, the MA powder is annealed at 600°C during which Sm 2Co 17 transforms to SmCo 5. The order of grain size in the annealed material is found to be less than 200 nm and the single domain particle size is of about 1 µm. As the coercivity strongly depends on grain size when it is less than the single domain particle size, MA provides hard magnetic powder at low temperatures without any sophisticated annealing programme. Conventionally, the preparation of SmCo5 magnets is based on high-temperature heat treatment above 800°C followed by a rapid cooling to room temperature. However, a suitable technique for consolidation of these MA powders is to be evolved. For MA soft magnetic material refer to Section 12.13. 12.9 MA POWDERS FOR SPRAY-COATINGS The MA process is suitable for the production of MCrAl (Y) (M=Fe, Ni or Co) types of powder like MA-754 superalloy which can be plasma-sprayed in a low pressure chamber to provide oxidation and corrosion-resistant coatings on gas turbine blades and vanes [43]. By employing a fine mixture of mechanically alloyed Cu/Al 2O3 composite powder for low pressure plasma spraying, a fine structure uniform coating has also been applied [44]. However, such coatings have been found consisting of two layers: matrix-rich layer and dispersoid layer. This separation presumably takes place due to the lack of wettability and to some extent by the acceleration caused in the plasma jet. A uniform single phase intermetallic coating of Ti 3Al has been obtained by employing MA Ti–Al powder for RF plasma spraying [44]. The interdiffusion of Ti/Al takes place quite quickly during the process of spraying, which results in the formation of Ti3 Al intermetallic phase. Moreover, an intermetallic matrix coating has also been obtained using MA Ti–Al–10vol.% Si 3N4 powder. A fine distribution of Si3N 4 is 129

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obtained along with the other phases like Ti5Si 3 and TiN exhibiting decomposition of the ceramic phase [44]. Using DC plasma spraying, a Cu/Ti/C MA powder coating of uniformly dispersed TiC in a copper matrix has been obtained [45]. These coatings have slightly higher hardness due to homogeneously dispersed structures. Thus, the proper selection of the ceramic phase both from the reactivity and wettability point of view is essential in order to have the uniform desired coating [44]. 12.10 SUPERPLASTICITY The MA 6000 oxide dispersion-strengthened, nickel-base superalloy has been found to possess superplasticity at 1000°C [46]. In the hotextruded and hot-rolled condition, i.e. prior to zone annealing, the material has a grain size of 0.26 µm and a dislocation density of 3×10 9 cm–2. It can be believed that at least some of the γ′ is evenly distributed as very fine (≈30 nm) particles (Fig.12.2, a similar microstructure, may be referred). The alloy exhibits superplasticity, having a maximum elongation of over 300% and a maximum strain rate sensitivity of 0.47. Transmission electron microscopy shows no evidence of ordinary recrystallization and grain growth is slight. However, strain rates of less than about 1.0 sec –1 alter the initial microstructure and prevent grain coarsening on subsequent annealing at higher temperatures. Deformation of the fine grained MA 6000 can be described as a combination of power law creep and diffusional (Coble) creep [47] (Fig.12.24). With a threshold stress caused by the presence of γ′ particles, it exhibits only the diffusional creep process. Threshold stresses for dislocation creep are not observed. Superplasticity has also been observed in MA IN 90211 aluminium alloy (Al–2%Mg–4.4%Cu–1.1%C–0.8%O) at 447°C and a strain rate of 1.0 sec –1. The alloy has a grain size of 0.5 µm. The deformation at a high strain rate is found to be governed by co-operative grain boundary sliding, i.e. the movement of grain groups as a unit [48]. Such a process proceeds as shear of grain groups along two intersecting systems, having grain boundary surfaces oriented at 45 to 60° with respect to the tensile axis. Superplasticity has also been observed in IN 9021-SiC p composite material (see Section 12.12). 12.11 TRIBOLOGICAL MATERIALS Aluminium alloys It is common practice to use solid lubricants such as graphite, tungsten disulphide or hexagonal boron nitride in sleeve-bearing materials to reduce friction in the ‘boundary lubrication region’. The MA process has proved to be a powerful tool in producing an ultrafine microstructure 130

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Fig.12.24 Comparison of the flow data for MA 6000 at 1000°C with the predicted dislocation and Coble creep behaviour of Ni (Ref.47, Met. Trans., 16A, 777 (1985) (The Minerals, Metals and Materials Society)).

containing dispersed solid lubricants such as graphite and tungsten disulphide [49,50], which has improved tribological properties. Al–4.5%Cu–graphite has been produced by MA of the elemental mixture in an attritor mill in an inert atmosphere for a period of 3 hrs followed by consolidation by hot pressing. The MA results in a homogeneous distribution of graphite particles in the matrix, which leads to a dramatic decrease in the friction coefficient in the boundary region (Fig.12.25) [49]. It is also interesting to note the slope of the curves in the boundary region, which is a measure of the rate of decrease in the friction coefficient f with respect to the ‘bearing parameter’ ZN/ P, where Z is the viscosity of the lubricating oil, N is the rpm of the shaft and P is the pressure on the bearing, which is faster in the case of MA materials. It is observed that an addition of approximately 3% graphite by MA is sufficient to improve the bearing performance of the alloy. Even in the presence of 3% graphite, the k-factor of the matrix remains unchanged, which is attributed to the refining of the grains and the strengthening effect of oxide and carbide dispersoids. The MA process also results in more than a 100% improvement in the wear resistance of the material. The MA technique has also been used successfully to improve the tribological performance of aluminium–lead alloys. Aluminium base-bearing alloys invariably contain tin as an alloying element which is costly. Lead is an attractive alternative to tin because it is much cheaper and more freely available. However, the Al–Pb system is immiscible with 131

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Fig.12.25 Coefficient of friction, f, vs bearing parameter, ZN/P.

a wide miscibility gap. In addition, the two metals are virtually insoluble in each other at room temperature. The situation is aggravated further by a large solidification range of their alloys and a wide density difference, which greatly increases the kinetics of lead segregation during melting and solidification. As a result of all these difficulties, it is virtually impossible to produce Al–Pb castings having a uniform distribution of lead. The MA of elemental aluminium and lead powders was carried out in an attritor mill for 2 hrs using 1% stearic acid as a PCA. After degassing at 200°C in a vacuum of 10 –2 torr, the MA powder was consolidated by hot pressing at 500°C under a pressure of 60 MPa [50]. During MA, lead particles undergo plastic deformation without becoming work-hardened as its recrystallization temperature (–3°C) is even below the ambient temperature. Thus, the lead particles are lamellar in shape and remain soft, whilst the aluminium particles which undergo plastic deformation, are work-hardened and become fragmented. Consequently, these fragments stick to the surface of the lead lamellae. The consolidated MA alloy has an improved microstructure, which is also reflected in an improvement of approximately 250% in the wear resistance of the alloy. High energy milling of Al, Fe and ferrochromium powders has been carried out to develop Al–8%Fe–2% Cr alloy. The reason for using chromium is that, being a carbide-forming element, it is likely to combine with the carbon of the PCA, therefore increasing the wear resistance of the alloy. After degassing, the powder was consolidated by hot pressing [51]. The MA material was found to possess a hardness of 162 BHN, which is comparable to several rapidly solidified Al–Fe–x-alloys. The MA alloy has high wear resistance compared to PM and wrought alloy, and EDAX analysis has shown the presence of Al4 (FeCr) in132

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termetallic compound. The alloy has the potential to be used in a wide variety of components for aerospace applications because of its high specific strength. Brasses The MA technique has also been used to improve the tribological properties of brasses. The MA of Cu–40%Zn blended alloy and solid lubricant WS 2 or graphite (20 and 40 vol.%) has been carried out in a planetary mill in an argon atmosphere for a period of approximately 5–25 hrs. It was observed that MA continuously refines the microstructure with milling hours (Fig.12.26) but in the case of brass WS2, milling beyond 13 hrs results in the chemical reaction of zinc with WS 2 to form ZnS due to activation by the energy input of MA. The graphite particles retain their hexagonal lattice, as confirmed by SAD patterns of the composite powder, even after milling for a period of 25 hrs. The MA powder containing WS 2 was consolidated by hot pressing at 620°C at a pressure of 28 MPa, and the powder containing graphite was consolidated by cold compaction under a pressure of 700 MPa followed by sintering at 810°C for one hour [52]. The microstructures of these consolidated MA composites have uniform fine scale distribution of WS2 graphite, as shown in Fig.12.27. Friction coefficient f (in unlubricated conditions) and wear rate measurements for these MA materials have shown that it could be possible to decrease the friction coefficient as low as that of the pure solid lubricants (f = 0.15), indicating that the sliding surface is completely covered by a thin layer of the solid lubricant.

Fig.12.26 TEM micrograph of Cu–40% Zn brass powder particle, milled for 13 hrs. Note the lattice planes of WS 2 (Ref.52, Metal Powder Report, 48 (11), 20 (1993) (Elsevier Science)). Fig.12.27 (right) Micrograph of sintered MA brass–graphite composite (Ref.52). 133

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Steels Wear resistant steels for tribosystems have also been developed using MA techniques as discussed in Section 12.2. 12.12 COMPOSITES Aluminium-base MMCs, having particulate reinforcements such as SiC, Al2O 3, spinel MgAl2O4 and BN, are being developed. However, a major problem with these composites is the excessive contamination by iron from the alloying mill abrasion. Unlike other processing routes for similar composites, the particulate reinforcement does not enhance significantly the strength of the matrix that has already been strengthened by the oxide and carbide dispersions. The matrix (IN 9021) is thermally unstable and loses its strength at temperatures above 150°C due to dislocation annihilation, and the coarse SiC particles do not appear to result in any significant strengthening at high temperatures [53]. Whilst the particle–matrix interface strength is substantial (well over 1000 MPa), coarse particles of SiC break during fast fracture and the fracture toughness of the composite is low [54]. However, at high temperatures in the range 425–450°C and at high strain rates (>0.7 sec –1), the composite material with the IN 9021 matrix shows extended ductility, up to 300% total elongation [53,55]. Elemental 7010 alloy has been mechanically alloyed with 10% SiC particles [56]. After degassing at 200°C for 2 hrs in vacuum (10 –2 torr), the composite is consolidated by hot pressing at 565°C under a pressure of 110 MPa and hot extrusion at 585°C with an 11:1 extrusion ratio. The consolidated composite was solution treated at 500°C and quenched in ice-water followed by ageing at 175°C for 16 hrs. The composite had a clean SiCP–matrix interface (Fig.12.28) showing minimum interface reaction. It should be noted that SiC P appears to inhibit the alloying

Fig.12.28 TEM micrograph showing SiCp interface in 7010 matrix (courtesy Prof. P. Ramakrishnan). 134

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process. It is also revealed that Zn dissolves with greater difficulty than Cu and Mg [57]. The matrix in the consolidated composite had dynamically recrystallized nanograins (20–50 nm) with a fine dispersion of oxides and carbides (5–15 nm) prior to recrystallization, resulting in highstrength properties in the composite material. The composite had a tensile strength of 100 MPa, almost double that of the matrix at 350°C. A composite with a 2014 aluminium alloy matrix having 8% SiC P has also been prepared successfully using the MA technique. The composite powder was consolidated by hot pressing at 535°C under a pressure of 125 MPa. It was observed that the MA process increases wear resistance of the composite by a factor of two in comparison with the conventionally blended composite [58]. Aerospace Metal Composites (AMC), UK has developed a technique to produce a 2124-T4 aluminium alloy matrix with 17.8% SiC P using MA. Blending as well as MA is carried out in an argon atmosphere to control oxidation, maintain product purity and reduce the explosion risk. Consolidation of the MA powder is carried out by HIP to produce a billet which is extruded into tubes to make MMC frames of road racing bicycles (Fig.12.29) [59] Al 2O 3 and spinel MgAl 2 O 4 have also been tried as the reinforcement with Al–4%Mg blended alloy matrix using the MA technique, resulting in composites with improved creep rates [60]. The spinel MgAl2O 4 has a stable structure up to 2000°C and is shown to minimize metal/ ceramic reaction. The MA aluminium-based, copper-based and iron-based tribo-composites were discussed previously in Section 12.11. In situ TiNi–TiC and Ni–TiC composites were prepared by MA of elemental powders. The powder was found to contain granulates of 310 nm after 3–5 hrs of MA due to a melt and quench mechanism taking

Fig.12.29 Raleigh's Cronos-MMC frame made of 2124-T4 aluminium alloy - 17.8% SiCp composite (XMT-100). 135

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place during the process. Spherical TiC grains with martensite and B2 structures also form during processing, which causes self-propagating synthesis initiated by the heat of formation of TiC [61]. Surprisingly, the MA process has been applied to polymer-based composites. The MA of powders of polytetrafluoroethylene and copper or nickel has been carried out. Amorphization through the SSAR has been achieved in the composite [62]. During MA, in general the softer material tends to form the matrix and the harder material disperses within it. This tendency, together with the tendency of mechanically alloyed materials to become harder with increasing processing, can be utilized to produce complex, multi-level metal composites. For example, if tungsten, a relatively ductile metal, is mechanically alloyed with a very fine ZrO2 nickel powder, it results in the dispersion of ZrO 2 in tungsten. If nickel powder is added and processed with tungsten ZrO 2 composite powder, a two-level composite is formed. The hard and brittle tungsten–ZrO 2 is broken up and dispersed in the matrix of more ductile nickel. Contact between the zirconium oxide and nickel is minimal in such a composite. Thus, a number of levels and the relative degree of dispersion of different ingredients that can be obtained by this technique are almost limitless. 12.13 AMORPHOUS SOLIDS Using the MA/MG technique, a large number of binary, ternary and even quarternary systems have been amorphized. A typical list can be seen in Table 5-1. Here, the few systems which have practical significance are discussed. At present, there is considerable interest in the formation of amorphous phases in soft magnetic alloy systems because of their potential properties and processing advantages. These materials have traditionally been prepared by the RS of molten alloys. Fe–B and Fe–B– Si containing 75% of the amorphous phase was prepared by MA [63]. Fe–Si alloy is used for the construction of transformers due to a combination of low magnetostriction and high saturation magnetization. Amorphous Fe–Si alloys, potentially, add to these characteristics the possibility of controlling magnetic anisotropies by stress and/or magnetic field annealing, facilitated by the atomic mobility of the amorphous state. The MA technique has been successfully used to amorphize the alloy [64]. Aluminium–transition metal (TM)–rare-earth (RE), quarternary systems (Al-TM1-TM2-RE) have yielded the best glass-formers to date. Optimal glass-forming compositions found in the Al–Ni–Fe–Gd system are located around Al 85Ni 6Fe3 Gd6 and Al85Ni 5Fe 2Gd 8. These amorphous alloys have a strength of 1280 MPa, which is 2-3 times higher than 136

Applications

the tensile strength of commercially available crystalline aluminium alloys. They are comparable to the strength of common high-strength steels which have low densities and an elastic modulus comparable with standard aluminium alloys, and have satisfactory corrosion resistance. Al 80Ni 8Fe 4Gd 8 alloy powder has recently been produced in the amorphous state by the MA technique (Spex mill, 160 hrs), which suggests the possibility of forming the bulk material through the powder metallurgical MA technique [65]. The MA amorphous powder was found to have good stability with a crystallization temperature of 620–650 KHN. K c is found to be in the range 1–5 MPa√m, identifying the produced powder as the brittle type amorphous phase. Dispersed ductile nanocrystals are found to improve the fracture toughness at the expense of hardness. The Ti–Si system is of particular interest for semiconductor devices, and Ti–Si intermetallic compounds such as Ti5 Si3 and TiSi2 are considered as high-temperature structural materials with extremely low specific weight, favouring their utilization in aerospace applications. The system has been amorphized by a planetary ball mill in an argon atmosphere [66]. In the case of a Ti 67.5 Si 32.5 powder mixture, an almost amorphous phase forms when milling takes place with intervals and continuous milling results in the Ti 5 Si 3 stable phase. In the case of Ti 33.3Si 66.7, the metastable C49 structure is formed. Both intermetallic phases exhibit a crystalline size typically between 10 and 15 nm. 12.14 NANOCRYSTALLINE MATERIALS In general, nanocrystalline solids are exclusively prepared by condensation from the vapour phase to a fine powder or by the sol–gel method (ceramics). In these materials, 50% of the actual volume consists of grain boundaries. Because of this large number of crystalline interfaces, an important fraction of the material has a disordered microstructure with no short-range order similar to a gas phase. As a result, nanocrystalline solids exhibit physical and chemical properties differing from those found in normal crystalline or amorphous materials, e.g. the diffusion coefficient in these materials can be as high as 10 9 times of the volume diffusion in a single crystal [67], which leads to enhanced solid solubility (even in the immiscible system with positive enthalpy of mixing) and, in certain cases, an increase of the catalytic activity. Carrying out MA for 40 hrs, the Mo solubility in the nickel lattice, which is approximately 8 at.% at room temperature equilibrium limit, could be increased by a factor of 3 (24 at.%) with an fcc nanocrystalline structure. The average particles size was, as studied by high-resolution transmission electron microscopy and image processing, 4 nm. The solubility limit was found to increase with a decrease in crystallite size. 137

Mechanical Alloying

The electrocatalytic activity, for the hydrogen evolution reaction in alkaline solutions, of the cold-pressed MA NiMo nanocrystalline cathodes is among the highest ever reported for this type of alloy [68]. This high activity appears to be directly related to the size of the crystallites [69]. It is reported that elemental Nb and Al powders have been mechanically alloyed for 5hrs in an argon atmosphere to synthesize nanocrystalline Nb 3Al. The MA powder has an average particle size of 4 nm. The mixture becomes two-phase, Nb 2Al (19%) and Nb3Al (81%), after heat treating for 1 hr at 600°C but has a stable grain size (5 nm). Grain growth is found to occur above 900°C [70]. Cu–16 at.% Fe and Fe–40 at.% Cu have also been produced in nanograin size using the MA technique [71]. Fe–50 at.% V alloy powder has also been nanosized (9 nm) by ball milling for 180 hrs [72]. High-temperature alloys such as Al–Ti, Ti–Al, Ti–Si and Nb–Al (T m ≈ 1850°C) have also been prepared in nanosize using MG with the aim of improving their mechanical properties as discussed in Sections 12.3 and 12.5. Reaction milling in the case of titanium (as discussed in Section 12.5) and Fe–Al [73] has been employed for nanocrystallization. 12.15 MECHANICALLY ACTIVATED CHEMICAL REACTIONS As discussed in Section 5, MA can be used effectively to reduce the T crit and a concomitant increase in reaction rate. These processes were applied to reduce a number of metal oxides and halides to pure metals [74]. Typically, copper metal forms when CuO is milled with calcium, and brass forms when CuO and ZnO are simultaneously milled with calcium. A break-through in the process came with the advent of a separating and extracting technique for the nanosized powder particles produced [75]. The nanosized particles have good chemical purity and surface oxidation less than 10–15%. So far, iron, nickel, cobalt and copper powders, as well as some ceramic powders such as alumina and zirconia, having a particle size as small as 5 nm, could be produced using the MA technique. These ultrafine powders have significant potential for high technology applications such as cutting tools, advanced ceramics, high density magnetic fluids and catalysts. Thus, the MA process allows direct production of metal powder without the need to manufacture bulk alloy and then convert it to powder form. A number of high-temperature processes can be combined into one single room temperature process with the potential for significant cost savings. It is envisaged that the process may find application in the production of reactive elements and alloys such as rare-earth and titanium, which are difficult to produce by conventional metallurgical processing techniques that require recycling of a process reductant. 138

Applications

Recently, using a uni-ball mill for milling a mixture of ilmenite and carbon, it was possible to produce Ti powder at a low temperature reduction of rutile [76]. The mineral ilmenite is a naturally occurring iron titanate (nominally FeTiO3 ) and is abundant in nature. Commercial grades of ilmenite contain 45–65.8% TiO2 and are regarded as a huge resource for the production of rutile (TiO 2), which can be used directly as a pigment or for the manufacture of titanium). The reduction rate was found to increase with milling hours and milling intensity. However, the carbothermic reduction does not take place during milling (even after 400 hrs of milling), but during subsequent low-temperature annealing at 760°C. Under normal conditions, this reduction usually takes place at temperatures between 860–1000°C. Thus, high-energy ball milling produces a mechanical activation effect, as shown in Fig.12.30. The gaseous reduction rate is enhanced due to the small grain/particle size and lattice defects created in ilmenite crystals, enabling complete ilmenite reduction of rutile to occur before an appreciable further reduction to TiO 2-x (where 2 < x < 0) occurs. Thus, MA leads to a high degree of selectivity also for the reduction of rutile. The effect of MA on elemental mixtures of Ti, Ni and C powders has also been investigated [77]. It is observed that an explosive reaction takes place during milling in the range 3 hrs 30 min–3 hrs 35 min, and large agglomerates with a size of 3–10 mm form abruptly, suggesting melting and subsequent quenching of the powders during MA. The phases present are indentified as spherical TiC grains, lath twin martensite, B2 phase and Ni depending on the initial composition. The self-sustained, high-temperature synthesis (SHS) is believed to be triggered by the release of the heat of formation of TiC and ignited by mechanical collisions. Thus, MA can be a versatile method for inducing SHS reactions in highly exothermic systems.

Fig.12.30 XRD patterns for the: a) 200 hrs milled sample isothermal annealed at 760°C for 30 min, and b) for an as-mixed sample heated continuously up to 1100°C without premilling, i; FeTiO 3 , r; TiO 2, F; Fe, A; γ-Fe (C) (austenite) (Ref.76). 139

Mechanical Alloying

12.16 OTHERS Other potential applications include dental prostheses [78] and superconductors [79–81]. The MA process is being used to produce precursor powders by INCO alloys based on Ba–Cu and Bi–Sr–Ca–Cu alloys for producing superconducting wires [82]. The use of ceramics for superconducting applications is not suitable because they are very brittle. CONCLUSIONS At present, MA materials are used not only in the aerospace industry, but also in the thermal processing industry (furnace fixtures, muffle tubes, furnace racks, furnace transport rollers, skid rails, mesh belts, electrical resistance windings, burner nozzles, etc.), the glass industry (molten glass stirring rods, furnace hearths, tiles, tableware, bushing used to produce single and multistrand fibres), the energy production industry (flame stabilisers, fuel cladding in fast neutron breeder reactors, heat exchanger components in high-temperature gas-cooled reactors, components for industrial gas turbines), the recreational industry (bicycle frames and fork brakes) and the spray-coating industry. However, the progress of industrial acceptance of MA materials has been slow, mainly due to the following reasons: – high cost (though high performance/cost ratio), – reluctance to try new materials, – lack of in-service experience, – non-availability of the necessary product forms (bar, plate, sheet, wire, tube, forging stock, etc.), – size range; the upper limit usually depends on the casting production facilities (at present 500 kg MA 956 plates represent the largest made), whereas the lower limit (thickness or diameter) may depend on the working characteristics of the alloy. The outlook for the growth of MA materials in industrial markets appears to be a good one. Efforts to increase the availability of the product forms and lower cost production methods will certainly improve the market scenario of these materials. The RSMA technique will probably result in the advent of a new class of extraordinary materials. References 1. 2. 3. 4. 5. 6. 7. 8.

R. Sundaresan and F.H. Froes, J.Metals., 39 (1), 22 (1987). R.C. Benn, L.R. Curwick and G.A.J. Hack, Powder Metal., (4), 191 (1981). K. Murakami et al, Met. Trans., 24A, 1049 (1993). W. Sha and H.K.D.H. Bhadeshiam Met. and Mater. Trans., 25A, 705 (1994). J.J. Fischer and J.H. Weber, Adv. Mat. Proc., (10), 43 (1990). T.S. Chou and H.K.D.H. Bhadeshia, Met. Trans., 24A, (4), 773 (1993). B. Kazimierzak, M. Prigon and D. Coutsouradis, MPR, 45 (10), 699 (1990). E. Kohler, C. Gutsfeld and F. Thummler, Powd. Met. Int., 22 (3), 11 (1990). 140

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C. Gutsfeld and F. Thummler, MPR, 45 (11), 769 (1990). J.F. Faure and L. Ackermann, Comparative evaluation of aluminium base materials and processing routes for elevated temperature applications, Presented at the symposium on Dispersion Strengthened Aluminium Alloys, TMS, Phoenix, Arizona (1988). R.D. Schelleng, J. Metals., 41 (1), 32 (1989). J.R. Pickens, J. Mat. Sci., 10, 1437 (1981). A. Layyous, S. Nadiv and I.J. Lin, Powd. Met. Int., 19 (1), 11 (1987). A. Layyous, S. Nadiv and I.J. Lin, Powd. Met. Int., 22 (3), 22 (1990). S.J. Donachie and P.S. Gilman, In: Aluminium-lithium alloys II, T.H. Sanders and E.A. Strake (eds.), TMS, Pennsylvania, PA (1984), p.507 R.D. Schelleng, P.S. Gilman and S.J. Donachie, In: SAMPE technical conference series (1985), p.106. R. Sundaresan and F.H. Froes, MPR., 44 (3), 195 (1989). P.K. Mirchandani and R.C. Benn, In: Spage Age Metals Technology, V.2, SAMPE (1988), p.188. K.S. Lee, S.J. Kim and K.B. Shin, In: Advances in Powder Metallurgy and Particulate Materials, V.4, A. Lawley and Swanson (compilers), MPIF and APMI, Princeton, N.J., USA. S.S. Ezz, A. Lawley and M.J. Koczak, Dispersion strengthened Al-Fe-Ce, a dual rapid solidification-mechanical alloying approach, Presented at the symposium on Dispersion Strengthened Aluminium Alloys, TMS, Phoenix, Arizona (1988). H. Baker, ASM Metals Handbook, V.2, Metals Park, OH (1979), p.269. A.N. Patel and S. Diamond, Mat. Sci. and Eng., 98, 329, (1988). W. Erde, PM Technology, 14, (3), 175 (1966). M.L. Mehta et al, Powd. Met. Int., 22 (5), 15 (1990). T. Morooka, J. Jpn. Soc. Powd. Metall., 41 (3), 277 (1994). J.L. Estrada, Mechanical alloying of prealloyed copper-lead powders, Presented at the 1996 World Congress on Powder Metallurgy, Washington DC (1996). J. Palacios, Production of copper-carbon and copper-aluminium-carbon alloys by mechanical alloying, Presented at the 1996 World Congress on Powder metallurgy, Washington DC (1996). R. Sundaresan and F.H. Fores, In: Proc. Int. Conf. PM and Related HighTemperature Materials, P. Ramakrishnan (ed), Oxford and IBH (1987), p.215. R. Sundaresan and F.H. Froes, MPR, 44 (3), 206 (1989). N.Hoo et al, J. Metals, 48 (7) 40 (1996). S.C. Glade amd N.N. Thadhani, Met. and Mater. Trans., 26A, 2565 (1995). T.D. Chan and C.C. Koch, Nanostruc. Mater., 5 (6), 615 (1995). E.Ivanov et al, J. Less Common Metals., 132, 25 (1987). K. Uenishi et al, Mat. Sci. & Eng., 134A, 1342 (1991). T. Yamane, Mater. Trans. JIM, 37 (2), 130 (1996). M. Mori, J. Jpn. Soc. Powd. Powd. Met., 38 (1), 71 (1991). W. Heisz and L. Schultz, Appl. Phys. Lett., 53 (4), 342 (1988). L. Schultz, K. Schnitzke and J. Wecker, Appl. Phys. Lett., 56, 868 (1990). K. Schnitzke, et al, Appl. Phys. Lett., 57 (26), 2853 (1990). M.Umemoto, J. Jpn. Soc. Powd. Powd. Met., 38 (7), 957 (1991). M. Ito, J. Jpn. Inst. Met., 59 (6), 666 (1995). J. Wecker, M.Katter and L. Schultz, J. Appl. Phys., 69 (8), 6058 (1991). P.S. Gilman and J.S. Benjamin, Ann. Rev. Mater. Sci., 13, 279 (1983). F. Fukumoto and I. Okane, In: Proc. Int. Conf. Thermal Spray Conference and Exposition, Orlando, Florida (1992), p.595. M. Umemoto and I. Okane, J. Jpn Inst. Met., 58 (1), 43 (1994). 141

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R. Singer and G.H. Gessinger, In: Proc. Int. Symp. Deformation of Polycrystals, N. Hansen et al (eds), Riso National Laboratory, Roskilde (1981), p.365. J.K. Gregory, J.C. Gibeling and W.D. Nix, Met. Trans., 16A, 777 (1985). M.C. Zelin, T.R. Bieler and A.K. Mukherjee, Met. Trans., 24A, 1208 (1993). A. Sharma, P.R. Soni and T.V. Rajan, MPR 43 (1), 37 (1988). V. Kumar et al, Unpublished work. G.H. Borhani, A.N. Tiwari and P. Ramakrishnan, Trans. PMAI., 20, 7 (1993). B. Günther et al, MPR., 48 (11), 20 (1993). Wadsworth, In: High-Strength Powder Metallurgy Aluminium Alloys II, G.J. Hildeman and M.J. Koczak (eds), TMS, Warrendale, PA (1986), p.137. D.L. Davidson, Met. Trans., 18A, 2115 (1987). R.G. Nieh, Scripta Metall., 18, 1405 (1984). A. Bhaduri et al, Trans. PMAI, 21, 1 (1994). A.P. Miodownik, Met. and Mater. Trans., 27 (10), 3718 (1996). N.R. Rudrappa et al, PMAI Newsletter, 11 (4), 23 (1985). S. Pickering, MPR, 50 (6), 30 (1995). J.W. Pyun and S.I. Kwun, J. Korean Inst. Metals,33 (6), 814 (1995). H.Q. Ye, Acta Mater., 44 (5), 1781 (1996). T. Ishida and S. Tamasu, J. Mater. Sci. Lett., 12 (23), 1851 (1993). S. Surimach, et al, Mat. Sci. and Eng., A134, 1368 (1991). A.G. Escorial et al, Mat. Sci. and Eng., A134, 1394 (1991). G.M. Dougherty, G.J. Shiflet and S.J. Poon, Acta Metall. Mater., 47 (2), 2275 (1994). Z.H. Yan, M. Oehring and R. Bormann, J. Appl. Phys., 72 (6), 2478 (1992). Y. Minamino et al, Met. Trans. JIM, 37 (2), 130 (1996). M.L. Trudean and R. Schultz, Mat. Sci. and Eng., A134, 1361 (1991). J. Kuyama, J. Jpn. Appl. Phys. Lett., 30 (5a), 854 (1991). Y. Tanaka, Mater. Trans. JIM, 37 (3), 265 (1996). J.Y. Huang et al, Acta Metall. and Mater., 45, (1), 113 (1997). H. Kuwano and T. Hamaguchi, J. Jpn. Soc. Powd. Powd. Met., 39 (3), 216 (1992). J. Rawers et al, Met. Mater. Trans., 27A, 3126 (1996). G.B. Schaffer and P.G. McCormick, Mat. Sci. Forum, 88-90, 779 (1990). P.G. McCormick, MPR, 51 (9), 5 (1996). Y. Chen et al, Met. and Mat. Trans., 28A (5), 1115 (1997). J.Y. Huang et al, Acta Metall. and Mater., 44 (5), 1781 (1996). E. Ivanov, Int. symp. Mechanical Alloying, Kyoto, Japan (1991). C. Politis, Physica, 135B, 286 (1985). M. Stubicar et al, Int. symp. mechanical alloying, Kyoto, Japan (1991). V. Mizutani et al, Mat. Sci. Forum, 88-90, 415 (1992). Anonymous, MPR, 43 (10), 698 (1988).

Questions 1. Give a general flow sheet for the production MA aluminium alloys. 2. What are DISPAL alloys? How are they produced? 3. Discuss development of MA high temperature aluminium alloys. 4. Discuss with the help of a flow sheet the production of Al-Li MA alloys. 5. Name the MA alloys produced at commercial level, discuss their properties and uses: (i)Nickel-base superalloys; 142

Applications

6. 7. 8. 9.

10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20.

21. 22.

(ii)Aluminium-base alloys. Which MA steels are produced on an industrial scale? What are their applications? What are ODMs? What are their possible uses? Discuss MA steels useful for tribological purposes. How MA techniques can be helpful in developing the following: (i)High temperature copper alloys; (ii)Ti-Mg alloys. How MG is helpful in improving ductility of titanium aluminades. Discuss how Ti 5Si 5 can be improved with the help of MG. How MA produces supercorroding materials. Discuss development of MA hydrogen-storage materials. Compare SSE achieved by MA with that by RS. How MA technique is helpful in developing Nd–Fe–B magnets. Write a note on MA powders useful for spray coatings. Why MA techniques are helpful in achieving superplastic behaviour. Give an account of development of tribological materials using MA techniques. Give an account of development of particulate composites using MA techniques. What are the advantages of nanocrystalline materials? Discuss one example where nanocrystallization has been achieved using MA techniques. Discuss how MA techniques are advantageous in production of titanium powder from ilmenite. List the reasons for the slow acceptance of MA products at industrial level. How can this state of affairs can improved?

143

Mechanical Alloying

144

List of Symbols Applications LIST OF SYMBOLS Aw Effective surface created during welding events; the weld area b Burger’s vector C Bmin Minimum concentration of solute atoms B Cp Specific heat of powder D Diffusivity D A , D B Self-diffusion coefficient of atoms A and B, respectively. d Particle diameter; grain diameter d0 Initial grain diameter E Young’s modulus Eb Energy dissipated during an impact E0 Pre-collision kinetic energy of the grinding medium ball Et Total energy dissipated during MA f Shape factor; coefficient of friction fw Force acting at the particles weld f(v)dv Fraction of particles having volume between v and v+dv Final shape factor f sf f si Initial shape factor G Shear modulus g Acceleration due to gravity g p , g r ,gt Geometrical constant with respect to Hertzian-impact pressure, radius and time, respectively H Powder hardness Hk Knoop hardness Hv Vickers hardness ∆H mix Enthalpy of mixing Ht Total enthalpy input by cold work h Height h Rebound height h0 Powder coating thickness ht Heat transfer coefficient I Intensity of a superlattice line k Particle thermal conductivity; pre-collection velocity constant with respect to impact time k0 Grain size coefficient km Materials constant kr Dislocation relaxation parameter with respect to applied tensile stresses l Dispersoid particles spacing N Revolutions per minute (rpm) n Number of rolling cycles; work hardening coefficient n(v)(dv) Number of particles having volume between v and v+dv Q Activation energy R Universal gas constant Rb Radius of the grinding ball Rc Radius of the mill-cell R ia i th peak position of the pair distribution function of the amorphous solid R ci Radius of the i th co-ordination sphere of the crystal R al, Rcl Nearest neighbour distance in amorphous and crystalline states respectively Rw Radius of weld region r Radius of atom r A ,r B Radius of atoms A and B, respectively 145

Mechanical Alloying rd rh rp S T Ta Tb Tc T crit Ti T ig Ts ∆T t tc t ig V Vp x xc Z αf αw ε e& ε c, ε f η ηb λ λ0 ν0 νr νs ρ σ σ th σu σy τ ω ωc ωp ωv

Dispersoid radius Hertz radius Particles radius Relative long range parameter Absolute temperature Ambient temperature Bulk temperature Curie temperature Critical temperature Post-impact temperature Ignition temperature Powder surface temperature Change in temperature Mechanical alloying processing time Time between collisions Incubation time for combustion to start Volume Volume of powder impacted per collision Vial length in Spex mill Critical thickness of amorphous layer Zener– Holloman parameter; viscosity of lubricant Particle fracture probability Particles welding probability Deformation manifested strains Particles strain rate during collision Critical strain to fracture Relative change in bond length Rebound yield Mean free path between collisions Atomic size factor Relative colliding velocity of the ball Grain boundary velocity during recrystallization Velocity of sound Density of ball Stresses experienced by powder during collision Threshold stress below which there is negligible creep Tensile strength (ultimate) Yield strength (0.2 per offset) Half of impact duration Angular velocity of horizontal ball mill Critical rotational velocity of horizontal ball mill Angular velocity of plate in a planetary mill Angular velocity of vial with respect to plate in a planetary mill

146

Index Applications

Index A acids 31 Ag–Cu 27, 125 Ag–Fe 27 Al–(7.5–15)% Ti 117 Al–2% Li–4.4% Mg 115 Al–2%Mg–4.4%Cu–1.1%C–0.8%O 130 Al–3% Li–2.2% Mg 114 Al–4% Mg–1.1% C–0.8% 112 Al–4.5%Cu–graphite 131 Al–8% Fe–4% Ce 118 Al–8%Fe–2% Cr 132 Al–Al4C3 composite powder 25 Al–Fe 27 Al–Fe–Ce 28, 116, 118 Al–Fe–Mn 28 Al–Mg 31 Al–Ti 116 Al2O3–Y2O3 105 Al50Mg50 33 aluminium alloys (see dispersion strengthened) amorphization 21 amorphization by MA 41 amorphization reaction 48 amorphous solids 136 Arrhenius equation 58 atomistic models 72, 87 attrition 6 attritor 6 attritor efficiency 11 attritor mill 3 average flake size 19

B B2 TiAl 26 ball mill 2 ball rotational velocity 81 ball to powder ratio 9 ball velocity gradient 11 ball–to–powder ratio (BPR) 19, 20 brazing 93 brittle–brittle system 37 bilk temperature rise 85

C CALPHAD 43 147

CALPHAD method 72 carbide dispersion 65 chemical reactions 51, 133 chromium carbides 102 (Co–Fe)75Si15B10 32 closed packed array 12 Coble creep 130 consolidation 58 co-operative grain boundary sliding (CGBS) 128 critical reaction temperature 52 cryomilling 26 crystal-to-amorphous transformation 47 Cu–Ag 29 Cu–Al 42 Cu–Al–C 116 Cu–C 116 Cu–40%Zn 133 Cu–8% Ti–4% B 120 Cu–Cr 119 Cu–Fe 4, 27 Cu–Pb 4 Cu–V 42 Cu–W 42 Cu–Zr 119

D dead zone 11 degree of crystallization 20 differential thermal calorimetry (DTC) 49 diffusion welding 92 diffusivity 53, 45 Dispal 91, 112, 4, 115 dispersion strengthening (DS) 2 materials 2 commercial materials 2 superalloys 1, 2, 3, 18, 100, 105 iron-base materials 4 aluminium alloys 107 copper 115 dispersoid particle radius 67 dispersoid radius 67 dispersoid spacing 67 dispersoid strengthening 67 dispersoid volume fraction 67 double mechanical alloying 28 ductile–brittle system 36

Mechanical Alloying

ductile–ductile system 35 dynamic compaction techniques 60

E EDAX 128 elongated grains 61, 62 equation of kneading 27 EPMA 39, 41 Energy maps 55, 56 ethylene bis disteramide 31 exothermic redox reactions 51 explosive compaction 89 extrusion ratio 60, 101, 113

F fatty acids 2 Fe–13.5%Cr–0.3%Mo–1%Ti–0.3%Y2O3 107 Fe–20% Cr–4.5% Al–0.5% Ti–0.5% Y2O3 107 Fe2B 22, 23 FeCrAl 109 fluid dies 59 fussion–fusion 77 forged bonding 94 friction coefficient 133 fracture dynamic 72 forging 72,74 shear 72 free-ball velocity 14 forging 60,62

G gas gun launcher 60 Ge–Si 37 glass transition temperature 43 glide 50, 66 global modelling 73 mechanistic 71 atomistic 71, 86 thermodynamic 71, 87 kinetic 71, 87 grind limit 17 grinding balls 9, 23 grinding media 10

H Hall–Petch relationship 65 Hall-Petch effect 67 head-on collisions 14 heat treatment 25, 28, 105, 106, 123 Hertz impact theory 84 Hertz radius 14, 84 Hertzian collision 85 Hf–Al 42 Hf–Cu 42 Hf–Ni 42 high-angle grain boundaries 49 high-energy milling mode 15 high-speed blenders and shakers 7 horizontal mills 81 hot isostatic pressing (HIP) 60 hot upsetting 104 HSLA steels 67 hydrogen storage materials 4, 123 hydrostatic stress 39

I ignition surface coating technique 2 ignition temperature 52 ilmenite 139 immiscible liquids 4 immiscible solids 4 impact energy 20 impact stress 14 impact time 83 IN 9021-T4 113 IN 9052 112 IN-905 XL 116 IncoMAP 111 INCONEL MA 754 103 incubation period 52 initial grain diameter 58 internal oxidation 2 International Nickel Company 3 interparticle necking 37 interparticle spacing 2 iron titanate 139

K ketones 31 kerosene 2 kinetic models 72

148

Index Applications kneading technique 27

MSMA 100, 101

L

N

laboratory planetary mill 9 large diameter ball mills 9 (see also horizontal mills) lattice strain 46 lattice defects 15 liquid phase sintering 61 liquid quenching model 41 local modelling 73 localized melting 61 long-range order (also see LRO parameter) low-angle grain boundaries 50 low density alloys 113 low energy milling mode 15 low pressure plasma spraying (LPPS) 125 LRO parameter 47

nanocrystalline cermets 51 nanocrystalline solids 137 nanocrystalline Ti–N powder 123 nanocrystallization 49 Nb3Ge 48 Nb–Al 42 Nb–Cu–Si 42 Nb–Ni 40, 42 Nb–Sn 48 NbC 110 Nb40Ni60 40 Nd–Fe–B magnets 126 Nd16Fe76B8 128 Ni–Zr 21 Ni–Zr2 48 Ni–20%Cr–2%ThO2 40 Ni32Ti–Ni45Nb55 46 Ni60Nb40 23 Ni60Ti40 23 nickel-base superalloys 3, 103 NiTi2–NiNb 46 NiZr2 22 NiZr2–Ni11Zr9 37

M MA 6000 alloy 91, 92, 103, 106 MA 754 103, 107 MA 758 103 MA 760 103, 104, 107 MA 956 107 MA 957 107 magnesium-base materials 123 magnetic materials 126 mean free path between collisions 80 mechanically activated chemical reactions 138 mechanical grinding 15 mechanical interlocking 58 mechanistic model 85 mechanistic models 72 metastable hydrides 25 Meyer’s hardness coefficient 68 MgAl2O4 134 MgO 112 Mg–Zn 16 Mg70Zn30 15 microcracks 38 microhardness 18, 29 mill parameters 19 mill speed 19, 21 mixing technique 2 modelling mechanical alloying 72 modified attritor 11 morphology 20, 37, 51, 101 Mössbauer spectroscopy 19 MS 100, 101

O occluding air 5 ordering 56 Orowan mechanism 66 Orowan strengthening 69 Orowan stress 70 Orowan–Ashby expression 67 Ostwald ripening 1 oxidation 1 oxidation resistance 103 oxide coated balls 10 oxide dispersion microforged material (ODM) 109, 110 oxide dispersion 65 oxygen/carbon ratio 32

P packing factor 81 pair distribution function 86 particle thermal conductivity 87 particle welding 15 pebble mill 3 pipe diffusion 36 planetary ball mill 9 149

Mechanical Alloying plasma short activated sintering 61 polytetrafluoroethylene (PTFE) 136 post impact temperature 87 potential energy of a ball 55 powder cooling 86 powder hardness 74 powder heating 85 precollision velocity 83 process control agents 31 processing time 80

R radial ball velocity profile 11 rapid omnidirectional compaction (ROM) 59, 60 rapid solidification 28, 96 reaction milling 25 rebound yield 55 relative long-range order 47 relative velocity of the balls at impact 74 repeated powder forging 29 repeated rapid solidification 41 repeated rolling 26 retarded temperature effect 22 RSMA process 100 rutile 139

S SSAR (see solid state amorphization reaction) SAD 133 SAP 1 SCC resistance 69 selective reduction process 2 self-sustained high-performance synthesis (SHS) 139 size of the grinding ball 20 Sm–Co magnets 128 Sm–Fe magnets 127 Sm–Fe–N 128 Sm-Co magnets 128 Sm-Fe magnets 127 Sm10.5Co89.5 129 Sm16.7Co83.3 129 sol–gel method 137 solid solubility extension 51 solid solution strengthening 2, 65 solid state amorphization reaction model 42 solid state welding techniques 91 Spex 8000 mixer mill 122 SPEX mill 80 Spex shaker mills 9

Spex vibratory mill 8 strain hardening coefficient 68 superconducting transition temperature 19 superconductors 4 superplasticity 4,66,125,130 supersaturated solutions 4, 125 Szegvari attritor 6 surfactants 31

T thermomechanical processing 62 Ti–1% Al–8% V–5% Fe–1% Er 123 Ti–1% Al–8%V–5% Fe–1% Er 101 Ti–24% V–10% Cr–5% Er 123 Ti–24%V–10%Cr–5% Er 101 Ti–24%V–10%Cr–5%Er 101 Ti–Cu 42 Ti–Mg–Gd 117 Ti–Mg 121 Ti–N 122 Ti–Si 36,136 TiNi–TiC 135 TiC 110 time–temperature excursion 58 time between collisions 80 TiN 110 titanium 1 titanium aluminides 121 titanium oxide 1 titanium systems 98 total stored energy 22 transformation path (see also amorphization path) 22,23,48

U uni-ball mill 12

V vacancies 15 Vegard’s law 125 vial 81 vial length 82 vibratory ball mill 8

W water-jacketed milling chambers 19 wear 20 welding 89,91 wettability 129 WS2 133 150

Index Applications

X X-ray technique 39 X-ray line broadening 19 (see also X-ray technique)

Y Y2O3 13, 36, 102 Yield strength 64,65,67

Z Zener–Holloman parameter 62 zone annealing 62,100 Zn–Fe 42 Zn–Ni 42

151